Reformer Tubes

Engineering Failure Analysis xxx (2010) xxx–xxx

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Engineering Failure Analysis

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Damage characterization in two reformer heater tubes after nearly10 years of service at different operative and maintenance conditions

Antonello Alvino a,*, Daniela Lega a, Francesco Giacobbe b, Vittorio Mazzocchi c, Antonio Rinaldi c

a ISPESL, Dipartimento Tecnologie di Sicurezza. Laboratorio Chimico e Tecnologico dei Materiali, Laboratorio Metallografico, Via del Torraccio di Torrenova 7,00133 Roma, Italyb ISPESL, Dipartimento Territoriale di Messina, Via dei Mille 89/bis, 98123 Messina, Italyc ISPESL, Dipartimento Tecnologie di Sicurezza, Via Alessandria 220/e, 00198 Roma, Italy

a r t i c l e i n f o

Article history:Received 12 April 2010Received in revised form 14 June 2010Accepted 14 June 2010Available online xxxx

Keywords:CreepCarburizationAgingReformer furnaces

1350-6307/$ - see front matter � 2010 Elsevier Ltddoi:10.1016/j.engfailanal.2010.06.003

* Corresponding author. Tel.: +39 0697895126; faE-mail address: [email protected] (A. Alv

Please cite this article in press as: Alvino A etdifferent operative and maintenance condition

a b s t r a c t

Reformer furnaces are used in petrochemical industry to produce hydrogen. Their mostcritical components are radiant tubes, where extreme temperature and pressure conditionsrequired the deployment of high alloyed austenitic HP grade steels, owing to their superiorstrength to creep rupture and good corrosion resistance. Nevertheless, these high strengthalloys undergo damage when process conditions allow coke deposition and maintenanceprocedures are not carried properly.

Two radiant tubes, coming from two different plants and made of HP40Nb and HP40Nbmicroalloy steels respectively, are investigated here to highlight the endured damage after85,000–96,000 h of service at maximum temperature of �950 �C. Light optical microscopy(LOM) and scanning electron microscopy (SEM) analyses both confirmed significant agingin both tubes, as revealed by a substantial microstructural evolution consisting of phasestransformation and precipitation, as well as cavitation and incipient microcracking in thebulk. Mechanical testing demonstrated the considerable worsening of mechanical proper-ties following such a microstructural deterioration.

However, noticeable differences between the two tubes were discovered as far as typeand distribution of damage, and extent of the aged zone. The small differences in alloyscomposition seem not sufficient alone to explain such dissimilarity in behaviour because,in addition to creep damage, one tube showed also clear evidence of carburization. Insteadit is argued that differences in process conditions (i.e. temperature, pressure chemical com-position of process fluids and steam–hydrocarbon ratio) and cleaning maintenance for dec-oking (i.e. type and frequency) may have synergistically triggered different damagemechanisms responsible for the different form of degradation on each reformer tube.

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1. Introduction

Reformer furnaces are widely used in petrochemical industry for large scale production of hydrogen by catalytic reactionsthat convert the reactants (hydrocarbons and steam) in hydrogen and carbon dioxide according to the following reactions.

CnHm þ nH2O! nCOþ ðm=2þ nÞH2 ð1ÞCOþH2O! CO2 þH2 ð2Þ

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al. Damage characterization in two reformer heater tubes after nearly 10 years of service ats. Eng Fail Anal (2010), doi:10.1016/j.engfailanal.2010.06.003

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These endothermic reactions are conducted in radiant columns consisting of tubes filled with catalyst and continuouslyheated up to an operating temperature exceeding 1073 K. Improvements in process and catalyst have resulted in increasedproduction and efficiency over time, with best performance achieved for a steam to hydrocarbon ratio around 3:1 and forhigher operating temperature and pressure simultaneously [1].

Radiant tubes are the most critical components of a reformer furnace, as they are exposed to severe conditions duringservice. Because of the extreme temperature, high alloyed austenitic HP grade steels are required since they show highercreep rupture strength and good corrosion resistance. In particular, centrifugal cast tubes are most often employed materials.

These components are generally designed for a minimum lifetime of 100,000 h of service at operating temperature of1173 K. Nevertheless, premature failures of furnace columns are frequently observed as a result of a number of damage mech-anisms, such as creep, carburization, oxidation, thermal shock, and accidental overheating [2–11]. In addition, the same chem-ical reaction involved in the process may be responsible, under certain conditions, for the formation and deposition of cokeonto the internal surface of the catalyst tube, adding extra damage besides what previously cited. In some cases, large defor-mation of tubes during service may bring columns closer together, which reduces heat dissipation from the tube surfaces andleads to the formation of hot spots in the material. Such a degradation causes a macroscopic deformation in the cross sectionof tubes, especially in the lower portion of the columns subject to higher thermal gradients. In these zones a noticeable tubeovalization is indeed often observed. This is taken as a meaningful damage index and is subject to monitoring. Two radianttubes, from two different reforming plants, were indeed replaced following the last non-destructive inspection because theyshowed an undesired degree of deformation. They are examined here to characterize the endured damage and to assess ifsmall differences in alloy composition, process parameters (e.g. service temperature, pressure, chemical composition of fuelgas, etc.), and maintenance procedures could induce appreciable differences in terms of degree and/or type of damage.

2. Technical background

The catalytic tubes coming from the two different hydrogen plants are conventionally labelled as F01 and F02 in this pa-per for convenience sake. Both reforming furnaces are top fired, containing a large number of catalyst columns arranged in aradiant chamber. Heat for the endothermic reforming reaction is provided through firing from burners, arranged accordingto the design of the chamber, and is transferred through radiation. The main design and service data for the two plants arelisted in Table 1. Accordingly, the reforming plants show slight differences in terms of service conditions, chemical compo-sition of the feed gas, number and dimensions of radiant tubes, productivity and service life.

The cleaning maintenance of both furnaces F01 and F02 was carried out during refinery turnarounds and minor serviceshutdowns. During refinery turnarounds (every 4 years), reformer tubes were subject to complete replacing of the catalyst.Decoking was also carried out by steaming (36 h) or pig cleansing. The latter one is an extra-accurate cleaning method usingair, nitrogen, liquid chemicals and special cleansing devices that remove mechanically the encrusted coke, as well as anyalien substance in the column. During minor service shutdowns (every 12–24 months) tube cleaning was additionally per-formed by steaming (24–36 h), with partial replacing of the catalyst, when necessary. Steaming was also carried out occa-sionally during extraordinary maintenance whenever the unscheduled shutdown lasted longer than 6 days.

During the service period (8 and 9 years for furnace F01 and F02 respectively), one pig cleansing and 14 steaming pro-cesses were totalled on catalyst tubes of furnace F01, whereas reformer columns of furnace F02 were subjected to decoking14 times, by pig cleansing on two occasions and by steaming the remaining ones .

In 2008, during a service shutdown after 85,000 h, all 204 reformer tubes of furnace F01 were inspected inside the fireboxwith automatic testing devices over the whole length of the tube. The inspection system contained an eddy current defectdetection system, capable of penetrating the complete tube wall, in addition to a dual laser device for the simultaneousmeasurement of the outside tube diameter. The eddy current inspection showed acceptable and stable microstructural con-ditions in most of the columns, except for four tubes where some indications of irregularities were detected. The measures of

Table 1Design data of the two plants.

F01 F02

Typology Top firing Top firingNumber of burners 60 arranged on 5 rows 60: 36 inners and 24 outersFeed gas composition Light benzine sulphur-free (70%) + GPL sulphur-free (30%) Fuel gas (80%) + GPL or methane (20%)Number of catalyst tubes 204 tubes arranged on 4 rows 176 tubes arranged on 4 rowsTube dimensions OD 128,3 mm OD 124,2 mm

ID 101,6 mm ID 101,6 mmTube service pressure 30 bar 25 � 27 barMaximum tube design pressure 33 bar 32,7 barTube service temperature (top) 530 �C max 560 �CTube service temperature (bottom) max 860 �C max 880 �CMaximum tube service skin temperature 890 �C 890 �CMaximum tube design skin temperature 955 �C 950 �CH2 production 61.000 N m3/h 42.000 N m3/hOperating service life 85,000 h 96,000 h

Please cite this article in press as: Alvino A et al. Damage characterization in two reformer heater tubes after nearly 10 years of service atdifferent operative and maintenance conditions. Eng Fail Anal (2010), doi:10.1016/j.engfailanal.2010.06.003


A. Alvino et al. / Engineering Failure Analysis xxx (2010) xxx–xxx 3

the outside tube diameters also confirmed these findings. It was concluded that the material was losing most of its ductilitydue to the aging accumulated after so many years of service under harsh operative conditions and severe load cycles. There-fore, inspectors recommended replacement of the four most damaged catalyst tubes to minimize the risk of a column failurebefore the next shutdown scheduled for 2010.

As for the second case of furnace F02, the geometric deformation of its tubes had been monitored differently, via a quan-titative laser method called LOTIS by which the radiant column stability is assessed from diameter deformation of the tubes.The critical condition is reached when the diameter deformation exceeds 5–7% [3], which was the case for 18 tubes in thisfurnace – judged to be too heavily damaged. Even in this case, most of them were replaced as a precaution to prevent suddenshutdowns from extraordinary maintenance.

3. Materials and methods

Catalyst tubes from furnace F01 were made of centrifugal cast austenitic HP40Nb grade, while tubes of furnace F02 weremade of centrifugal cast austenitic HP40Nb + microalloy grade. Table 2 shows the chemical composition of the alloys. Tubesegments with 1.5 m in length were cut from both furnace regions and were sent to the laboratory for characterization bymeans of metallographic, chemical and mechanical analyses. The exact locations where the sections were cut on the refor-mer columns in the respective plants were unreported and remain unknown.

Each tube portion, as-received, was further severed into two blocks (labelled arbitrarily upper and lower) to facilitate col-lection of metallographic samples from each block. Moreover, in order to make a comparison between different service condi-tions, it was decided to compare the materials against the as-cast, pre-service conditions. Unfortunately, a direct experimentalcomparison could be done only for material F02. To this end, additional samples were purveyed from a third tube, named F0, inas-cast conditions and with chemical composition nearly identical to tube F02. Instead, because no samples in pre-service con-ditions were found for tube F01, a comparison was carried out against the literature data for as-cast HP–Nb steel [12].

All samples in this study were subject to metallographic analyses using optical microscope (LOM) and scanning electronmicroscope (SEM) operated in backscattered mode. Metallographic specimens were taken from each tube to examine bothtransverse and longitudinal sections. The cutting was performed by a precision saw (Buehler Isomet 4000). Specimens weremounted in an automatic mounting press (Buehler Simplimet 1000) and carefully polished using an automatic grinder-pol-isher system (Buehler Phoenix Beta) using a series of grit papers and finished by diamond polishing down to 1 lm. After-wards, polished areas were etched with the Murakami’s reagent (10 g K3Fe(CN)6, 10 g KOH, 100 ml H2O) or Groesbeck’sreagent (4 g KMnO4, 4 g NaOH, 100 ml H2O) and examined first under a Nikon Eclipse ME600 optical microscope. Some sam-ples, serving as controls, were examined in un-etched conditions too. Then, SEM analysis was carried out by a field emissiongun FEG-SEM (Zeiss LEO 35) equipped with backscattered electrons detector and microanalysis for Energy Dispersive X-rayanalysis (EDX). The Vickers hardness of each tube samples was measured by a durometer (Wolpert) according to the relatedUNI EN ISO 6507-1 standard [13]. At last, destructive tensile tests were performed on standardized specimens machinedfrom longitudinal blank cuts. The tensile tests were carried out at room temperature in a 60 ton capacity test machine(W + B), according to the related UNI EN 10002-1 standard [14].

4. Results

4.1. Visual examination

Pictures of the as-received tubes for the investigation are shown in Fig. 1. There is no obvious evidence of wall thinning orcracks at a visual inspection. Tube F01 exhibits an inner surface massively affected by coke scaling. A cross section of tubeF01 highlights this feature as a brown ring adherent to the inner surface and visible by eyesight. On the contrary, the innerwall of tube F02 was smooth and clean, without any trace of corrosion or scaling. These differences may be traced back eitherto cleaning and maintenance operations, or to the chemistry of the reforming process, since local deposition of coke is knownto occur if the steam/hydrocarbon ratio is not adjusted to about 3:1 and excess of hydrocarbon is present.

4.2. Microstructure

4.2.1. Light optical microscopySamples from tube F0 (the as-cast pre-service condition corresponding to F02) showed an almost identical pattern to that

reported by De Almeida and co-workers [12,15], namely: a typical network of primary carbides within a dendritic pattern

Table 2Chemical composition of the analyzed alloys.

C Si Mn Cr Ni Nb Ti Cu Fe

F01 0.40 1.50 1.50 25.00 35.00 1.50 – – Bal.F02 0.43 1.40 1.00 25.00 35.00 1.50 Var.a 0.25 Bal.

a Proprietary addition: % Ti ranges from 0.05 to 0.30.



Fig. 1. Photographs of the reformer tubes sent for investigation.

Fig. 2. Micrographs of cross sections of tube F0 at different magnifications.


lying in a austenitic matrix that is free of precipitates, as confirmed from observation at higher magnifications (Fig. 2a and b).A similar structure should be also expected for the material F01 in as-cast conditions, the only difference being a more con-tinuous skeletal pattern of primary carbides due to the lack of titanium [12].

A very different trend was observed in samples from serviced material. The LOM analysis of tube F01 (Fig. 3) put in evi-dence interdendritic carbides, strongly coarsened and forming a continuous network spanning the entire cross section of thetube. A massive precipitation of secondary carbides was also clearly visible as a dispersion of particles within the austeniticmatrix, with the only exception of the portion of metal close to the outer surface where precipitates are less abundant andlarge cavities are often detected. A progressive increase in microstructural degradation of the bulk material from the uppersection to the lower portion of the radiant tube was observed (data not shown), such that specimens from the latter partexhibited the largest degree of coarsening of dendritic carbides and the most conspicuous secondary carbides precipitation.

Near the inner surface of the tube, the metal in the wall was more affected by microstructural evolution, as shown inFig. 4. The overview in Fig. 4 reveals three different zones. The first one (Z1 on the micrograph) is the farthest from the wallinner side and it is characterized by coarsened dendritic carbides and by some secondary carbides within the austenitic ma-trix, which is still observed. Closer to the inner surface (i.e. moving from left to right on the micrograph), a second zone (Z2)extended for about 400 lm is encountered, showing very coarsened interdendritic precipitates and a massive secondary pre-cipitation so that the austenitic matrix is hard to be distinguished. The metal closest to the wall inner side (Z3 on the micro-graph) has an even different appearance, since this third zone extends radially for about 500 lm from the inner tube surface(which is exposed also to process fluids) and presents a particular structure with a large number of longitudinal and radialcracks. It is worth noting that in some samples this area is less extended or even almost absent, probably due to the brittlenature of its constituents that makes this portion of the material easy to detach and export away either in service or duringthe cleaning maintenance of the catalyst tube.



Fig. 3. LOM micrographs of cross sections of samples from tube F01: (a) near outer surface, (b) near inner surface, (c) bulk and (d) higher magnification of b.

Fig. 4. LOM micrograph of a sample from tube F01 (cross section), showing three zones of different microstructural evolution and damage, increasing fromZ1 (bulk, on the left) to Z3 (the damaged zone close to inner surface, on the right).


Samples from tube F02 analyzed by LOM showed an appreciable degree of coarsening of the primary carbides, though notof the same degree of tube F01 (Fig. 5). The continuity of the interdendritic carbide network seems to match the interden-dritic regions of the original cast structure. With respect to tube F01, the carbide network is less continuous. Finer and moreevenly distributed secondary precipitates are present within the austenite matrix, in agreement with De Almeida et al. [12],especially in the portion of tube near the wall inner side. All the analyzed specimens show a similar microstructural evolu-tion, since no appreciable difference was detected in either the degree of coarsening or the density of precipitates betweensamples from upper and lower portion of the tube (data not shown). A restricted zone adjacent to the wall inner side ap-peared to be almost free of precipitates and no indications could be found of the damage observed in the wall inner sideof tube F01. Nevertheless, some isolated voids and cavities were present in this area. A similar pattern is observable alsoclose to the wall outer side of F02, where large cavities are present.



Fig. 5. LOM micrographs of cross sections of samples from tube F02: (a) near outer wall side, (b) near inner surface, (c) bulk and (d) higher magnification ofc.

Fig. 6. SEM BSE images illustrating cross sections of samples from tube F0 at different magnifications.


4.2.2. SEM analysisSEM analysis on samples from tube F0 using backscattered electrons rendered the same microstructure observed by LOM

but revealed that the interdendritic phase consists of two different precipitates, which can be recognized by brightness con-trast in Fig. 6. The EDX microanalysis indicated that the dark phase is chromium-rich, while the white regions are precipi-tates rich in niobium and titanium (Fig. 7). This brightness contrast is due to the difference in atomic weight of chromium vs.niobium. The same can be expected for tube F01 in the as-cast condition [12].

The SEM analysis with backscattered electrons of un-etched specimens from tube F01 provided more detailed informa-tion about the patterns previously described from LOM microscopy. The region near the outer surface showed niobium-richsecondary phases (Fig. 8a) but not chromium carbides, likely because of massive chromium depletion and outward diffusionrelated to the development of a chromium-rich oxide layer [16], as suggested by EDX analysis (Fig. 9). Large, isolated cavitieswere also apparent, presumably due to the Kirkendall effect [17].

Approaching the wall inner side, a progressive increase in density and coarsening (Fig. 8b–e) of interdendritic precipitatescould be appreciated in the bulk material. EDX microanalysis revealed that there are three different phases that can be dis-tinguished from each other on a brightness basis in Fig. 10. The dark second phase is a chromium-rich carbide with smallamount of iron and nickel, whereas a white region, which is very similar in appearance to the bright phase observed inthe as-cast material, contains significant amounts of nickel, silicon and niobium, along with slight quantities of iron and



Fig. 7. EDX spectra of sample from tube F0 illustrated in Fig. 6b: (a) chromium-rich dark grey precipitates and (b) bright precipitates. The spotted areas areindicated by arrows in Fig. 6b.

Fig. 8. SEM BSE images illustrating cross sections of samples from tube F01: (a) near outer surface, (b) bulk, (c) higher magnification of b, (d) bulk, withmore massive precipitation, (e) the same as d, but at higher magnification and (f) near inner surface.


chromium. Consistent changes in the concentration of these elements have been observed by microanalysis in selected areas.It is worth noting that SEM–EDX compositional analysis is semiquantitative in nature and the depth of the penetration is



Fig. 9. EDX spectra of sample from tube F01 shown in Fig. 8a. The spotted zone are evidenced by arrows and circles: (a) bright phase, (b) dark layer,identified as chromium oxide, (c) chromium-depleted austenitic matrix and (d) austenite with normal content of chromium.


generally larger than the precipitate size so that it is very difficult to measure exactly a single small particle. Since it has beenreported that niobium carbide is not stable in the range of temperature between 700 and 900 �C, transforming into a nickel-niobium silicide generally known as G-phase [18–22], these bright areas probably contain this intermetallic compound to-gether with the NbC carbide not yet transformed. The EDX analysis supported this hypothesis. Also, chromium-rich second-ary carbides were easily observed within the austenitic grains as dark grey particles. Isolated cavities and voids weredetected approximately in the two thirds of the tube wall section towards the inner surface (Fig. 11a–c). They were mainlyarranged at the dendritic boundaries of primary carbides, as priorly evidenced by Le May et al. [23], and to a minor extentinto the secondary carbides. In a few cases, small cracks have also been detected in this portion of material, originating fromboth interdendritic and secondary precipitates. Chromium-rich precipitates are more prone to form cavities and microcracks,than niobium-rich carbides and G-phase.

The tube portion nearest to the wall inner side showed a large number of longitudinal and radial cracks, as depicted inFigs. 8f and 11d–f. This zone corresponds to the aforementioned brown ring observable on the cross section of the tube andhas a totally different microstructure than what previously described for the bulk material. The chemical composition of thephases appeared also to vary, showing both chromium-rich and chromium-depleted austenite regions, as well as zones withappreciable amounts of oxygen, silicon iron, and nickel (Fig. 12). As observed by Swaminathan [11] and De Almeida [12],



Fig. 10. EDX spectra of sample from tube F01 shown in Fig. 8e. The spotted zone are evidenced by arrows and circles: (a) dark grey precipitates, (b) whiteregion identified as G-phase, (c) brightest zone in the circle, identified as residual NbC and (d) secondary precipitate, chromium-rich.


cracks propagate into the bulk material along the interdendritic precipitates, attacking preferentially chromium-rich car-bides. Instead niobium carbides and G-phase seem to be more resistant to this attack, as shown in Fig. 13 and are removedfrom the matrix probably via mechanical erosion only because EDX microanalysis of the dark, cracked zones revealed sig-nificant amounts of oxygen, it is thought that these zones are mainly constituted by carbides, metal oxides and productsof local corrosion. Further evidence of dramatic changes, both in chemical composition and microstructure, were also foundin its magnetic properties. It is well known that this HP steel, when not carburized, is paramagnetic, but it gradually shows aferromagnetic behaviour as the intensity of carburization increases, due to the loss of chromium from the austenitic matrixin form of chromium-carbide precipitation [24,25]. Chips machined from the inner wall side of tube F01 were checked andfound to be ferromagnetic, indicating that tube F01 has indeed suffered significant carburization.

A marked microstructural evolution was also observed from samples of tube F02 in terms of interdendritic phases coars-ening and fine precipitation of secondary carbides caused by service aging (Fig. 14). The appearance of the outer surface wassimilar to tube F01, with a chromium-depleted zone and presence of niobium-rich precipitates and large cavities. EDX anal-ysis revealed the presence of both chromium and titanium oxide forming a thin layer (Fig. 15a). The bulk material showedboth interdendritic and secondary precipitates. As rendered by EDX analysis, secondary precipitates are chromium-rich andappear as dark particles with either cubic or needle-shaped morphologies, in agreement with Ribeiro [26]. Analogously to



Fig. 11. SEM BSE images illustrating the microstructural evolution on samples from tube F01, as previously described in Fig. 4: (a) Z1 zone, (b) Z2 zone, (c)particular of microcrack in Z2 zone, (d) Z3 zone, (e) and (f) particular of voids and cracks in Z3 zone. Cross sections.


tube F01, the interdendritic phases showed different brightness at the EDX microanalysis (Fig. 16) depending on their chem-ical composition, with dark areas being chromium carbides, and brighter spots being most abundant in niobium, nickel andsilicon. Significant amounts of titanium were sometimes detected at greyer regions surrounding the brighter precipitates(Fig. 16d), which suggests a partial conversion of (Nb–Ti)C to G-phase and titanium carbide, although not in the same mas-sive form observed in the case of tube F01 that did not contain titanium. The wall inner side appeared almost free of pre-cipitates (with the exception of rare niobium-rich phases) and was covered with a brown film that was confirmed to bechromium oxide by EDX microanalysis (Fig. 15b). A large number of voids and cavities is visible throughout the entire sec-tion of these samples, especially near the wall inner side. In contrast to what were reported for tube F01, both niobium-richphases and chromium carbides appeared to be equally affected by cavities and cracks in our material (Fig. 17).

4.3. Mechanical strength properties

4.3.1. Hardness valuesVickers hardness tests were performed on each reformer tube sample at 30 kg load. Both the upper and lower blocks of

tubes were examined. Each block was divided in four sectors, namely outer wall, bulk-1, bulk-2 and inner wall. Six readingson each sector were taken, in order to detect, if any, an increase in hardness due to the massive carbide precipitation. Themean values from these readings are reported for each specimen in Table 3. The hardness value of material in as-cast con-dition (F0) is 179 ± 7. As expected, there is a correlation between the change in hardness and the amount of carbide precip-itation [21]. In fact, tube F01, which shows massive carbide precipitation, rendered appreciably higher values than the as-cast material, with highest readings near the inner side of the tube wall (that is the tube portion most affected by precipi-tation). On the contrary, no difference between as-cast and aged material were found between tube F02 and its counterpartF0, as well as no appreciable difference in hardness values could be detected along their cross section.



Fig. 12. EDX spectra of samples from tube F01 in Z3 zone: (a) bright zones; (b)–(d) dark areas and crack zones.


4.3.2. Tensile strength propertiesTensile specimens of standard size were machined out from longitudinal blanks cut from both upper and lower blocks of

the tube samples supplied. Tests were carried out in a 60 ton capacity test machine at room temperature. Since data of theas-received tube samples were not available, the tensile tests were compared with that supplied from the manufacturer forthe as-cast material [27,28] in Table 4. Notably, tensile strength values obtained for aged materials of both tubes F01 and F02were remarkably lower than the as-cast material. Only ultimate tensile strengths are reported because it was not possible todetect the 0.2% proof strength due to a brittle failure, without any appreciable elongation and necking. The examination offracture surfaces with a stereomicroscope confirmed the brittle nature of fracture. The reduced tensile ductility of the tubesamples validate that exposure to service conditions had degraded the mechanical properties to an unacceptable level.

5. Discussion

Both serviced materials analyzed in this study parade a substantial degree of aging, as put in evidence by the dramaticloss of mechanical properties in association with a strong microstructural evolution characterized by transformation andprecipitation of phases, cavitation, and incipient local microcracking in the bulk material. For both serviced tubes, the degreeof microstructure evolution and damage increases approaching the wall inner side.



Fig. 13. SEM BSE micrographs of samples from tube F01, dealing the preferential dissolution of chromium-rich carbides. Niobium-rich phases (G-phase andNbC) in the carburizing zone appear as white isles surrounded by voids, cracks, and corrosion products. Some of them have been evidenced by circles.

Fig. 14. SEM BSE micrographs of samples from tube F02: (a) outer surface, (b) bulk material, (c) higher magnification of bulk and (d) inner surface. Crosssections.


Anyways, significant differences in damage exist between tube F01 and F02. On the one hand, the alloy F01 showsextended coarsening of interdendritic precipitates, massive secondary precipitation of chromium carbides and an almostquantitative interconversion of niobium carbide into G-phase. Moreover, cavitation and cracks propagate along dendriticboundaries by dissolving preferentially chromium-rich precipitates. On the other hand, in the alloy F02 secondary precip-itates are finer, the conversion of NbC into G-phase seems to be less massive than in tube F01, and cracks affect bothniobium-rich and chromium carbides in the same degree. As for the damage distribution along the tube, tube F01 doesnot show a uniform extent of aging from the upper to the lower portion. Even though carburization is likely to be active



Fig. 15. EDX spectra of samples from tube F02 illustrating the different chemical composition of the oxide layers at inner and outer surface: (a) chromiumand titanium oxide layer at outer surface and (b) chromium oxide film at inner surface.


over the entire tube length with the same degree, a visible damage gradient is observed, with a progressive increase inmicrostructural degradation of the bulk material from the upper section to the lower portion of the radiant tube. Spec-imens from the latter part exhibit the largest degree of coarsening of dendritic carbides and the most conspicuous sec-ondary carbides precipitation. In our opinion, these observations suggest that the lower section of tube F01 sufferedgreater aging, perhaps being a hotter spot. On the other hand, all specimens from tube F02 show approximately the samemicrostructural evolution, pointing to a uniform aging process along the entire length of the portion analyzed of this cat-alyst tube.

Another noticeable difference in the microstructural evolution between the two tubes resides in the metal degradation atthe inner side of the radiant tube wall. The microstructure of tube F01 closely resembles that reported in literature for car-burized HP steels in either cracking or reforming plants [24,29]. It can be assumed that significant carbon diffusion had oc-curred, leading to partial carburization of the steel at the inner surface of the tube as a consequence of the exposition tocarbon-rich process fluids. The carburized material lost its alloying elements in the form of carbides, becoming susceptibleto corrosion [30,31]. This could explain the large number of cracks clearly observable in this tube portion. As mentioned,these cracks propagate into the bulk material along the interdendritic chromium-rich carbides network and these zonesare mainly constituted by carbides, metal oxides and products of local corrosion according to EDX analysis. The inner surfaceof tube F02 is different, being covered by a thin layer of chromium oxide that protected the material from metal degradationand carburization. Yet, this zone had suffered creep damage, as revealed by extended cavitation.

The loss of mechanical properties of tube F01, proved by very low tensile strengths, is compatible with the carburizationargument reported in literature [32]. Vickers hardness test results also agree with this hypothesis, since the values measurednear the wall inner side of tube F01 are appreciably higher than those collected in the bulk, pointing out a hardness increaseinduced by the additional precipitation of carbides. Further evidence of the significant carburization suffered by the materialF01 is provided by the cited ferromagnetic behaviour of the metal near the inner wall side due to the loss of chromium fromthe austenitic matrix.

Collectively taken, these findings indicate that the difference between tubes F01 and F02 in either the extent of aging orthe type of damage can be ascribed to the specific process conditions of the two reformer furnaces, which are not identical(Table 1) and to the maintenance procedures followed during service shutdowns. In particular, we offer that the chemicalcomposition of the process fluid, and primarily its correct balance with steam, played a key role in reducing lifetime ofsuch components, since the excess of hydrocarbon is a well known common cause for coke deposition [33] and, hence,for the progressive decrease in heat transfer to the process fluid, which in turn reduces process yield and promotes car-burization and creep damage. With all probability the different decoking history of the furnaces played a crucial role ondamage evolution. Steaming was performed in furnaces at least once a year, whereas the accurate tube decoking by pigcleansing was carried out two times on F02 and only one time on F01. Since coke removal by steaming is usually done attemperatures higher than the design operational limit, detrimental effects on tube life may occur. So, relatively frequentdecoking may possibly have accelerated creep damage, as verified by SEM on both tubes F01 and F02, while – in addition– the less frequent pig cleansing may have steered a more massive coke deposition on tube F01, causing its partialcarburization.



Fig. 17. SEM BSE micrographs of samples from tube F02 evidencing the cracks propagating across both chromium carbides and niobium-rich phases. Crosssections.

Fig. 16. EDX spectra of sample from tube F02 shown in Fig. 13c. The spotted zone are evidenced by arrows and circles: (a) brown chromium-richprecipitate, (b) bright niobium-rich particles evidenced by the circle, identified as niobium carbide, (c) bright precipitate, identified as G-phase and (d)titanium-rich precipitates near niobium-rich particles.




Table 3Vickers hardness values from test results performed on the supplied materials.

Sample Upper Lower

Tube Outer wall Bulk 1 Bulk 2 Inner wall Outer wall Bulk 1 Bulk 2 Inner wall

F01 211 ± 8 217 ± 9 216 ± 8 243 ± 10 215 ± 6 212 ± 3 217 ± 7 252 ± 5F02 172 ± 10 170 ± 10 167 ± 10 181 ± 5 172 ± 10 163 ± 7 162 ± 6 175 ± 10

Table 4Ultimate tensile strength values from test results performed at room temperatureon the supplied materials. For reference materials, data supplied by manufacturerhave been used [27,28].

Sample Upper LowerTube

F01 294 ± 15 267 ± 20F02 320 ± 22 339 ± 10

ReferenceG 4852 470G 4852 micro 470


6. Conclusions

We studied the damage and microstructural evolution of two radiant tubes which had been in service in reforming plantsfor about 10 years. The difference in material performances, aging degree and forms of damage was pointed out by a numberof techniques. The significant degradation of mechanical properties was observed in both tubes, which lost most of their duc-tility according to tensile test results. However, the most severe degradation in tube F01 entailed an appreciable increase ofhardness directly correlated to microstructural features of the material.

Optical and electronic microscopy allowed to ascertaining structural evolution of both serviced alloys, providing detailedinformation about the damage mechanisms involved in the degradation and aging of the tubes. Both tubes F01 and F02 havesuffered creep damage, since voids, cavities, and small cracks could be found in each case.

In agreement with mechanical tests, tube F01 displayed a greater degree of damage. Besides microstructural evolutionand creep damage, in facts, there was evidence of carburization inside tube F01. The results support that the small differ-ences in the chemical composition of the alloys, particularly the addition of Ti in tube F02, do not convincingly explain suchdissimilarity. Even though the titanium in tube F02 is admittedly beneficial for creep resistance and lifetime increase, themassive degradation observed on the inner surface of tube F01 could result from other factors, e.g. process temperatureand pressure, chemical composition of fuel gas, and especially the steam/hydrocarbon ratio (favouring coking and subse-quently carburization), which all can contribute to the onset of this form of damage. Nevertheless, despite small differencesin process conditions (Table 1) and material composition (Table 2) are present between furnace F01 and F02, the dramaticchanges in damage evolution observed in the two furnaces cannot be ascribed exclusively to these factors. Instead, it’s likelythat the process parameters have influenced the damage and aging of the analyzed components in synergy with the differentdecoking strategies (i.e. type and frequency) used in the two cases. So, the steaming process, which was frequently carriedout on both reformer furnaces, is likely to have accelerated the creep damage of both tubes, while the lower frequency ofdecoking by pig cleansing on furnace F01 may have played a key role in promote carburization on tube F01.

Acknowledgements

The authors are grateful to Linde Gas and Raffineria di Milazzo for providing catalyst tubes and technical support aboutreforming processes. The help of Mr. Antonio Di Loreto for samples preparation and Mr. Ubaldo Iglio for performing mechan-ical tests is gratefully acknowledged. Prof. Silvia Licoccia of University of Rome ‘‘Tor Vergata” for granting access to the FES-EM lab is also acknowledged.

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