A

tsbEoTs©

K

1

rCmmraTpd

atteSic

0d

Materials Science and Engineering A 460–461 (2007) 392–402

High temperature corrosion resistance of candidate nickel-basedweld overlay alloys in a low NOx environment

R.M. Deacon ∗, J.N. DuPont, A.R. MarderLehigh University, 5 East Packer Avenue, Bethlehem, PA 18015, USA

Received 5 January 2006; received in revised form 17 January 2007; accepted 19 January 2007

bstract

Changes in environmental regulations have led many fossil fuel-fired boiler operators to alter their combustion practices (low NOx burning),hereby lowering plant emissions. This change has led to unacceptable wastage of carbon and low alloy steel waterwall tubes and expensivehutdowns due to severe corrosion. One favored solution is to weld overlay a more corrosion resistant alloy on top of existing tubes. Two nickel-ased alloys developed for such applications were tested alongside the commercially available alloy 622 in a simulated low NOx environment.lectron probe microanalysis (EPMA) examination of the weld overlays and corrosion scales demonstrated that microsegregation of molybdenum

ccurred in one of the candidate alloys and alloy 622. This microsegregation had a detrimental effect on the corrosion resistance of these alloys.he candidate alloy with higher chromium concentration, low nominal molybdenum concentration, and corresponding minimum molybdenumegregation, exhibited the best corrosion resistance of the examined alloys. 2007 Elsevier B.V. All rights reserved.

gtAttass

hrbfte[n

eywords: Microsegregation; Nickel-based alloys; Sulfidation

. Introduction

In an effort to reduce boiler emissions in accordance withecently implemented environmental regulations, such as thelean Air Act, many coal-fired power plant operators haveoved toward a staged combustion process. By delaying theixing of fuel and oxygen, and thereby creating a reducing envi-

onment in the boiler, the amount of nitrous oxides (NOx) thatre released as a by-product of combustion is reduced [1,2].he use of staged combustion has been found by many powerlant operators to be the most cost and time effective method forecreasing NOx emissions.

Prior to implementation of staged combustion, most boilertmospheres were oxidizing, allowing for formation of protec-ive metal oxides on waterwall tubes made out of grades 11 or 12ype steels [3,1]. Under those conditions, failure due to accel-rated waterwall wastage was generally not a major problem.

taged combustion boilers, on the other hand, create a reduc-

ng atmosphere in the boiler due to the lack of oxygen. Sulfurompounds from the coal are transformed into highly corrosive

∗ Corresponding author. Tel.: +1 610 758 4270; fax: +1 610 758 6407.E-mail address: rmd3@lehigh.edu (R.M. Deacon).

atmps

l

921-5093/$ – see front matter © 2007 Elsevier B.V. All rights reserved.oi:10.1016/j.msea.2007.01.150

aseous H2S [4]. Subsequent reaction with the metal waterwallubes leads to the formation of metal sulfides on tube surfaces.dditionally, corrosive deposits may form on the waterwall

ubes due to the accumulation of solid particles in the combus-ion environment, such as ash and un-burnt coal. In the reducingtmosphere of the staged combustion boiler, low alloy steels areusceptible to excessive wastage of the tube and unsatisfactoryervice lifetimes [4,1].

One favored solution to the problem of waterwall wastageas been to deposit a weld overlay cladding of a more corrosionesistant alloy on to the tube. Commercially available nickel-ased superalloys, such as alloys 622 and 625, have been usedor weld overlays [5]. These alloys provided more protection inhe reducing environment than standard steels, but are consid-red expensive and are susceptible to circumferential cracking6]. Initially designed for strength at high temperature, theseickel-based alloys contain alloy additions that play little, ifny, role in improving the alloy’s corrosion resistance, but addo the cost. Considering the large surface area of tubing that

ust be protected, there is considerable demand in the fossil fuel

ower generation community for a weld overlay alloy designedpecifically for corrosion resistance in a low NOx environment.

In this program, two experimental nickel-based weld over-ay alloys developed by ThyssenKrupp VDM USA Inc. were

R.M. Deacon et al. / Materials Science and Engineering A 460–461 (2007) 392–402 393

Table 1Substrate and filler metal compositions for the alloys in this study (wt%)

Element Substrate Alloy 33 wire Alloy 50 wire Alloy 622 wire

Cr 0.30 (max) 32.85 19.55 21.11Ni 0.30 (max) 30.95 51.6 58.22Fe Balance 32.2 13.7 4.23Mo 0.08 (max) 1.67 11.55 13.53Co ND 0.16 0.75 0.77W ND 0.09 1.6 3.73Ti 0.03 (max) 0.01 0.01 NDNb 0.02 (max) <0.010 0.24 NDCu 0.30 (max) 0.58 0.11 0.06Mn 0.80–1.40 0.64 0.38 0.31Si 0.40 (max) 0.31 0.09 0.2N 0.012 (max) 0.39 0.111 NDP 0.025 (max) 0.01 0.009 0.027SC

cNbaaattlta

2

5ipiwosrma

TW

P

PSSVCPBTWI

Fa

(ctpc

0.015 (max) 0.004 0.002 <0.0010.20 (max) 0.007 0.005 0.012

orrosion tested in an environment modeled after a typical lowOx boiler atmosphere. The commercially available nickel-ased weld overlay alloy 622 was tested alongside the candidatelloys to serve as comparison. The corrosion product on eachlloy was examined and related to the weld metal compositionnd distribution of alloying elements in the coating. The objec-ive of this study was to develop a fundamental understanding ofhe high temperature corrosion behavior of candidate weld over-ay, and from this, identify cost effective experimental alloy(s)hat have improved corrosion resistance over currently usedlloys.

. Experimental procedure

Weld overlays of two experimental alloys (termed 33 and0) were prepared using an automated gas metal arc weld-ng (GMAW) process. Filler metal compositions and weldingarameters are shown in Tables 1 and 2, respectively (weld-ng parameters for the alloy 622 overlay used for comparisonere not available). Cross sections of the experimental weldverlays were mounted and prepared for metallographic analy-

is (Fig. 1). All-weld-metal samples for corrosion testing wereemoved from all three weld overlays by wire electro-dischargeachining (EDM) and ground to 600 grit SiC on all faces, edges

nd corners. The samples were immersed in 4% picral solution

able 2elding parameters for alloys 33 and 50

arameter Alloy 33 Alloy 50

osition Flat Flathield gas Cronigon He 30 S Cronigon He 30 Shield gas flow rate 18 l/min 18 l/minoltage 13.5 Vdc 30 Vdc

urrent 180 A 140 Aeak current/time 400 A/2.7 ms 400 A/2.7 msase current/frequency 50 A/90 Hz 50 A/98 Hzravel speed 23 cm/min 35 cm/minire feed speed 2.1 m/min 6 m/min

nter-pass temperature 60 ◦C 130 ◦C

m

ztpiw1watc

iHui

ig. 1. Cross section photomicrographs of weld overlays (a) alloy 33 and (b)lloy 50. LOM, electrolytic chromic acid etchant.

etchant) to check for the presence of base metal; because theorrosion resistance of only the weld metal was of interest inhis study, those testing positive were rejected. Corrosion sam-les were measured to the nearest hundredth of a millimeter,leaned in acetone and then weighed to the nearest tenth of ailligram on a digital balance.Gaseous corrosion testing was carried out in Lindberg hori-

ontal tube furnaces equipped with high purity alumina processubes, Kanthal heating elements and Platinel-II thermocou-les. Weld metal coupons were suspended in alumina cruciblesn order to collect any spalled corrosion scale. The samplesere heated at 50 ◦C/min and held at 500 ◦C for 100, 500,000 and 2000 h. The normalized weight gain of the couponas determined by measuring the coupon weight before and

fter the test. Measurements of the coupon weight after theest included any spalled material that was captured in therucible.

The gas used for the corrosion tests was modeled after a typ-

cal [7,8,9,10] low NOx environment: 10% CO–5% CO2–2%

2O–0.12% H2S–N2 (vol%). The gas was supplied by the man-facturer without any added water vapor. Water was introducednto the hot zone of the furnace, using a syringe pump with a

394 R.M. Deacon et al. / Materials Science and Engineering A 460–461 (2007) 392–402

Fsl

ct

cacfi

siwtpmwda

TWt

E

CNFW

Fig. 3. Secondary electron image of alloy 50, showing location of EMPA linescan (a) and composition of iron, nickel, chromium and molybdenum along thisline (b).

ig. 2. Secondary electron image of alloy 33, showing location of EMPA linecan (a) and composition of iron, nickel, chromium and molybdenum along thisine (b).

apillary tube. The flow rate of the pump was adjusted to providehe correct amount of water vapor.

To allow for comparison of the candidate test alloys to aommercially available weld overlay alloy, alloy 622 was testedlongside the candidate alloys. In the tube furnace tests, the twoandidate alloys and alloy 622 were tested simultaneously in oneurnace, to eliminate any variation in the data due to differencesn furnace conditions.

Corroded test coupons were mounted under vacuum in coldetting epoxy and ground through 600 SiC. Samples were pol-shed to 0.05 �m surface finish. Light optical microscopy (LOM)as used to analyze the surface and polished cross sections of

he corrosion scales. Qualitative analysis of corrosion scale com-ositions and quantitative microsegregation traces in the weldetal was acquired on a JEOL 733 Microprobe equipped with

avelength dispersive spectrometers. Wavelength and energyispersive spectroscopy was performed at an accelerating volt-ge of 20 keV and a phi(ρz) correction scheme was used to

able 3eld metal compositions (wt%) iron, nickel and chromium weld metal concen-

rations (wt%) and calculated weld dilutions for overlays used in this study

lement Alloy 33weld metal

Alloy 50weld metal

Alloy 622 weldmetal

r 27.89 17.98 NDi 28.23 49.5 52.29e 38.56 18.26 8.5eld dilution (%) 11.1 5.8 7.3

Fig. 4. Secondary electron image of alloy 622, showing location of EMPA linescan (a) and composition of iron, nickel, chromium and molybdenum along thisline (b).

R.M. Deacon et al. / Materials Science and Engineering A 460–461 (2007) 392–402 395

Ft

cd

3

3

tmzi

ttcca

ig. 5. Normalized weight gain data for 100, 500, 1000 and 2000 h exposureso the mixed oxidizing/sulfidizing in horizontal tube furnace.

orrect for absorption and fluorescence of X-rays that occursuring electron probe microanalysis (EPMA).

. Results and discussion

.1. Initial characterization of weld overlays

The typical welding microstructures of the alloys used inhis study, including the cellular or dendritic weld metal, the

artensitic structure in the base metal and the partially mixedone, are presented in Fig. 1. No welding defects were observedn the welds.

The dilution of the weld overlays was measured to ensurehat the overlays were representative of overlays produced in

he field, where dilution values are typically 5–15%. The con-entrations of iron, nickel and chromium in all-weld-metaloupons removed from the overlays were measured; only ironnd chromium were measured for alloy 622 (Table 3). The dilu-

Fig. 7. LOM cross section of alloys 33 (a), 50 h (b) and 622 h (c) after 2000 hcorrosion test.

Fig. 6. Macrographs of test coupons after exposure to mixed oxidizing/sulfidizing gas at 500 ◦C.

3 ce an

t

D

w(((

i

ooaTfc

96 R.M. Deacon et al. / Materials Scien

ion was calculated using Eq. (1):

= Cifz − Ci

fm

Cibm − Ci

fm

(1)

here Cifz is the concentration of element i in the fusion zone

Table 3), Ci the concentration of element i in the base metal

bmTable 1), Ci

fm the concentration of element i in the filler metalTable 1), and D is the weld dilution.

The dilutions for both candidate alloys were calculated forron, nickel and chromium. These three values were averaged to

wcl1

Fig. 8. Electron microprobe maps of alloy 33 coupon from 2000 h

d Engineering A 460–461 (2007) 392–402

btain an overall dilution for the weld sample. Iron and nickelnly were averaged for the dilution value of alloy 622, as themount of chromium in the 622 weld metal was not measured.able 3 lists the average dilutions for the three alloys; it wasound that the dilution values of the overlays in this study areonsistent with those observed in the field.

Quantitative compositional analysis was performed on the

eld overlays to measure the amount of microsegregation of

hromium, iron, nickel and molybdenum (Figs. 2–4). EPMAine scans were performed across several cells for each alloy at.5 �m step intervals. The distribution of alloying elements in

tube furnace test—entire corrosion product: 20 keV, 65 nA.

nd En

tcncdi(mamm

dagfscd

R.M. Deacon et al. / Materials Science a

he weld overlays was found to vary depending on the alloy, espe-ially for molybdenum, which is known to segregate strongly inickel-based alloys [11]. Alloy 33 (2 wt% molybdenum initialoncentration) exhibited very little microsegregation of molyb-enum (Fig. 1). The segregation of chromium, iron and nickeln this alloy is minimal as well. For alloys 50 (Fig. 3) and 622Fig. 4), the difference between the maximum and minimum

olybdenum concentration is approximately 7 wt% for both

lloys. During solidification of these alloys, microsegregation ofolybdenum ahead of the solidification front occurred, creatingolybdenum enriched interdendritic regions and molybdenum

aftu

Fig. 9. Electron microprobe maps of alloy 50 coupon from 2000 h

gineering A 460–461 (2007) 392–402 397

epleted dendrite cores. The low diffusivity of molybdenum inustenite [12] prevents back-diffusion down the concentrationradient, leading to the cored structure shown in Figs. 3 and 4.Tourther quantify the microsegregation in these alloys, it is pos-ible to calculate the equilibrium coefficient for iron, nickel,hromium and molybdenum for each alloy from the EPMAata. If dendrite tip curvature effects are ignored, and it is

ssumed that solid state diffusion is negligible, liquid state dif-usion is infinitely fast and that equilibrium is maintained athe solid/liquid interface, then the Scheil relationship may besed to determine the composition of any solid (Cs) during

tube furnace test—entire corrosion product: 20 keV, 65 nA.

398 R.M. Deacon et al. / Materials Science an

Table 4Calculated equilibrium distribution coefficients of alloys 33, 50 and 622

Iron Nickel Chromium Molybdenum

AAA

s

C

wso

C

ecdmdmdeda

eaceot

C

fiFaIEtfo1c

3

B2Il

A

o5ssodh

fosao

vwdaever, is compact and thin, indicating that it is more effectiveat preventing continued attack than the products formed onalloys 50 and 622. The presence of thick, nodule like growthson the surface of the alloy 50 and 622 coupons (Fig. 6)

lloy 33 1.07 1.00 0.94 0.59lloy 50 1.14 1.02 0.95 0.67lloy 622 1.21 1.22 0.82 0.72

olidification:

s = kCo(1 − fs)k−1 (2)

here k is the equilibrium coefficient, fs the fraction solid duringolidification and Co is the nominal composition. At the onsetf solidification when fs is 0, Eq. (2) can be simplified to:

s = kCo (3)

Determining the value of k is then simply a mathematicalxercise, wherein the values of Co are known from the alloyhemistry and the values of Cs are dendrite core compositionsetermined using the EPMA line scans. For chromium andolybdenum, which were shown to segregate to the interden-

ritic regions, the dendrite core compositions are taken as theinimum values from the line scans; for iron and nickel, the den-

rite core compositions correspond to the maximum values. Thequilibrium coefficients are summarized in Table 4. From thisata, it is obvious that molybdenum segregates most strongly inll three alloys, with k values much less than 1.

It should also be noted that while alloy 33 exhibited the low-st equilibrium coefficient for molybdenum (0.59 versus 0.67nd 0.72 for alloys 50 and 622, respectively), this does not indi-ate that alloy 33 would exhibit more severe segregation of thislement than the other two alloys. The minimum compositionf the solid is given by the following equation, again assuminghat Scheil conditions exist:

min = kCo (4)

Using the line scan data and the calculated equilibrium coef-cients, the values of Cmin can be calculated using this equation.or alloy 33, the minimum molybdenum concentration is 1 wt%nd the maximum concentration detected using EPMA is 2 wt%.n contrast, for alloy 622, the minimum concentration given byq. (4) is 9 wt%, while the maximum is 17 wt%. Thus, even

hough the equilibrium coefficient for molybdenum is smalleror alloy 33 than alloy for 622, the lower nominal concentrationf molybdenum in alloy 33 versus alloy 622 (1.7 wt% versus3.5 wt%) leads to a much smaller range of molybdenum con-entrations in alloy 33.

.2. Corrosion testing

The results of the corrosion exposures are presented in Fig. 5.y combining the data from the separate 100, 500, 1000 and000 tests, an overall kinetic curve can be plotted for each alloy.

t is evident from this figure that alloy 33 exhibited a markedlyower corrosion rate than both alloys 50 and 622.

Fig. 6 displays the alloy coupons after each exposure.lloys 50 and 622 developed friable scales, while the scale

Ff

d Engineering A 460–461 (2007) 392–402

n alloy 33 was less prone to spallation. The scale on alloy0 after 1000 h exposure is extensively spalled; although thecale remained relatively adherent to the coupon during expo-ure, large sections of the scale broke away during removalf the coupon from the test crucible. In contrast, alloy 33id not experience any significant spallation during couponandling.

The polished cross-section of the 2000 h exposure couponsor each alloy is presented in Fig. 7. The scale that developedn alloy 33 is much thinner (tens of microns) compared that thecales on alloys 50 and 622 (hundreds of microns). This figurelso illustrates the open, porous structure of the external scalen alloys 50 and 622.

The weight gain data correlates well with the optical obser-ations of the corrosion test coupons. Alloys 50 and 622,hich exhibited higher corrosion rates compared to alloy 33,eveloped thick corrosion products that were non-adherentnd very porous. The corrosion product on alloy 33, how-

ig. 10. Electron microprobe maps of inner scale region of alloy 50 couponrom 2000 h tube furnace test—entire corrosion product: 20 keV, 65 nA.

nd En

an

3

tdpamtp

ptptdmsrh

R.M. Deacon et al. / Materials Science a

lso indicates that the corrosion product on these alloys wason-protective.

.3. Qualitative corrosion product analysis

The corrosion products that formed on all three alloys inhis study were complex and consisted of several layers ofiffering composition. Several researchers have reported multi-hasic corrosion products forming on nickel-based alloys with

dditions of iron and chromium in sulfur containing environ-ents [13,14]. Electron probe microanalysis was performed on

he 2000 h corrosion test coupon for each alloy; the results areresented in Figs. 8–12. Due to the thickness of the corrosion

iac

Fig. 11. Electron microprobe maps of alloy 622 coupon from 2000

gineering A 460–461 (2007) 392–402 399

roducts on alloys 50 and 622, two maps were produced forhese alloys: one which includes the entire corrosion productlus some weld metal, and another higher magnification maphat focused on the internal scale only (internal corrosion here isefined as any corrosion product that forms beneath the originaletal surface). While they do not provide quantitative compo-

itional data, these maps are a valuable tool for determining theelative composition of the corrosion product and understandingow the corrosion product formed.

Comparison of the low magnification maps (Figs. 8, 9 and 11)llustrates the differences in corrosion product that form on eachlloy. On alloy 33 (Fig. 8), an external scale rich in iron, nickel,hromium, oxygen and sulfur developed. A thin (∼10 �m) and

h tube furnace test—inner corrosion product: 20 keV, 65 nA.

4 ce an

cnilo

isgwo

wcTdiws

00 R.M. Deacon et al. / Materials Scien

ontinuous chromium and oxygen rich layer formed in the inter-al corrosion product, very close to the interface between thenternal scale and the weld metal. It was also noted that veryittle molybdenum was incorporated in the corrosion product,wing to the low concentration of molybdenum in this alloy.

Alloy 50 (Fig. 9) developed an external scale that was richn iron, nickel and sulfur, with only a few areas in this outercale containing detectable amounts of chromium and oxy-

en. The inner scale is dominated by chromium and oxygen,ith some presence of molybdenum and sulfur. Examinationf the inner corrosion product only (Fig. 10) reveals regions

Fic

Fig. 12. Electron microprobe maps of inner scale region of alloy 622 coupon

d Engineering A 460–461 (2007) 392–402

ithin this chromium layer that are simultaneously depleted inhromium and oxygen, and enriched in molybdenum and sulfur.he microsegregation of molybdenum in the weld metal to inter-endritic regions (illustrated previously with EPMA line scans)s also evident in this figure. Note that the segregation from theeld metal was incorporated directly into the inner corrosion

cale as it formed.The low magnification map for alloy 622 is presented in

ig. 11. As with alloy 50, the external corrosion product is richn iron, nickel and sulfur, with an inner product that containedhromium, oxygen and molybdenum. The inner corrosion prod-

from 2000 h tube furnace test—inner corrosion product: 20 keV, 65 nA.

nd En

uda5

moplmiplaio

ptrdtcto

mnHdawtcom

nt5srisrlts

mtnsdmno

mTh

4

y

•

•

•

•

ddrmct

Acknowledgements

R.M. Deacon et al. / Materials Science a

ct (Fig. 12), containing mostly chromium and oxygen, is alsoecorated with areas that are depleted in chromium and oxygen,nd enriched in molybdenum and sulfur, as was seen with alloy0.

It is clear that the scales developed on alloys 50 and 622 werearkedly different than the scale on alloy 33. A chromium and

xygen rich layer formed near the weld metal/internal corrosionroduct interface for all three alloys. On alloy 33, however, thisayer is thin and continuous, while on alloys 50 and 622, it is

uch thicker and is decorated with regions that are depletedn chromium and oxygen. Additionally, the external corrosionroduct on alloy 33 is considerably thinner than the externalayer on alloys 50 and 622. This suggests that the thin chromiumnd oxygen rich layer in alloy 33 was more effective at prevent-ng continued attack of the weld metal than the chromium andxygen rich layer on both alloys 50 and 622.

While the presence of chromium and oxygen in a corrosionroduct does not necessarily signify the presence of the protec-ive oxide Cr2O3, it is likely that the chromium and oxygen richegions of the scale were more protective than areas that wereepleted of these two elements. It has been well documented inhe literature [15–17] that Cr2O3 is a protective oxide. Alloy 33ontained the most chromium of all the alloys in this investiga-ion, and this was an important factor in the corrosion resistancef this alloy.

It is also not possible to determine from these maps if theolybdenum and sulfur rich layers in the scale are molybde-

um sulfide; the stoichiometry of this scale was not determined.owever, molybdenum sulfide-type phases are known to beefective [18], and these regions were hypothesized to have acteds fast diffusion paths through the scale. Alloys 50 and 622,hich exhibited these molybdenum and sulfur rich regions in

he chromium and oxygen rich inner scale, exhibited excessiveorrosive attack. Alloy 33, which developed a more continu-us inner corrosion product with minimal presence of sulfur orolybdenum, appeared to be more corrosion resistant.The origin of the chromium/oxygen depleted, molybde-

um/sulfur enriched regions of the scale was determined to behe microsegregation that occurred upon solidification. Alloys0 and 622 exhibited marked segregation of molybdenum, ashown in Figs. 3 and 4. During exposure to the corrosive envi-onment, the corrosion front advancing into the weld metalncorporated the molybdenum enriched regions into the corro-ion product, resulting in the molybdenum and sulfur enrichedegions illustrated by the EPMA maps. Alloy 33 exhibited veryittle microsegregation of molybdenum in the weld metal, andherefore molybdenum-enriched regions did not develop in thecale that formed on this alloy.

Several investigators [18–20] have found that additions ofolybdenum to nickel and nickel chromium alloys decrease

he corrosion rate compared to similar alloys without molybde-um additions. However, homogenized alloys were used in thosetudies, in which there were no molybdenum concentration gra-

ients. Thus, the phenomenon observed here, in which regions ofolybdenum enrichment formed in the corrosion scale, would

ot occur in homogenized alloys. The corrosion scales that formn homogenized alloys are uniform and less likely to develop

nDa

gineering A 460–461 (2007) 392–402 401

olybdenum rich regions that could serve as fast diffusion paths.his explains why additions of molybdenum to the alloys studiedere had a negative effect on corrosion resistance.

. Conclusions

The results of the corrosion testing and compositional anal-sis can be summarized as follows:

Alloy 33 exhibited the best corrosion resistance of the threealloys examined in this study. This alloy contained the mostchromium and the least molybdenum, allowing for the forma-tion of a continuous chromium-rich inner corrosion productthat may have been responsible for the performance of thisalloy in simulated low NOx environments.The corrosion products that formed on the alloys in this studywere complex and not composed of a single compound. Thecorrosion scale that developed on the alloy depended on thecomposition of the overlay. Alloy 33 contained mostly iron,nickel and chromium, and the corrosion scale that formed onthis alloy contained those elements along with oxygen andsulfur. Alloys 50 and 622 contained significant molybdenumadditions, and this element was also found in the corrosionscales on these alloys.The presence of molybdenum in the weld metal served toundermine the corrosion resistance of alloys 50 and 622. Themicrosegregation of this element during solidification of theweld led to an inhomogeneous distribution of molybdenum inthe weld metal. The advancing internal corrosion front incor-porated this inhomogeneous distribution into the corrosionscale. The presence of molybdenum rich areas in the scalecaused the corrosion kinetics to increase due to their actionas high diffusivity paths for corroding species.Alloy 33, with less than 2 wt% molybdenum, did not experi-ence the microsegregation of this element observed for alloys50 and 622, and was therefore able to develop a more con-tinuous and uniform corrosion product. This corrosion scalewas more protective than the scales observed on alloys 50 and622 and led to the reduced kinetics observed for this alloy.

It was concluded from this work that additions of molyb-enum to alloys intended for use as weld overlay coatings isetrimental rather than beneficial to the corrosion resistance ineducing environments due to the microsegregation of this ele-ent. Weld overlay alloys should be designed to achieve their

orrosion resistance through alloy additions that are less suscep-ible to microsegregation, such as chromium or aluminum.

This work was supported by ThyssenKrupp VDM Tech-ologies Corporation, USA. The authors wish to acknowledgeave Ackland for assistance with the electron microprobe

nalyses.

4 ce an

R

[

[[

[[

[[[17] F.H. Stott, F.I. Wei, Oxid. Met. 31 (1989) 369–391.[18] M.F. Chen, D.L. Douglass, Oxid. Met. 32 (1989) 185–206.

02 R.M. Deacon et al. / Materials Scien

eferences

[1] C. Jones, Power (1997) 54–60.[2] R. Whitaker, EPRI J. (1982) 18–25.[3] J.A. Urich, E. Kramer, FACT (Am. Soc. Mech. Eng.) 21 (1996) 25–29.[4] S.C. Kung, W.T. Bakker, Mater. High Temp. 14 (1997) 175–182.[5] G.D. Smith, C.S. Tassen, Mater. Perf. 28 (1989) 41–43.[6] K. Luer, J.N. DuPont, A.R. Marder, C. Skelonis, Mater. High Temp. 18

(2001) 11–19.[7] J.E. Gabrielson, E.D. Kramer, ASME Joint Power Generat. Conf. 1 (1996)

19–23.[8] W.T. Bakker, J. Stringer, Mater. High Temp. 14 (1997) 101–108.[9] Y. Kawahara, Mater. High Temp. 14 (1997) 191–198.10] S.F. Chou, P.L. Daniel, High Temp. Corros. Energy Syst., Proc. Symp.

(1985).

[[

d Engineering A 460–461 (2007) 392–402

11] J.N. DuPont, Metal Trans. 27A (1996) 3612–3620.12] S.W. Banovic, J.N. DuPont, A.R. Marder, Sci. Tech. Weld. Join. 7 (2002)

374–383.13] H. Xu, M.G. Hocking, P.S. Sidky, Oxid. Met. 39 (1993) 371–388.14] C. Fang, H. Yakuwa, M. Miyasaka, T. Narita, Oxid. Met. 53 (2000)

597–617.15] F.H. Stott, G.C. Wood, J. Stringer, Oxid. Met. 44 (1995) 113–145.16] H. Hindam, D.P. Whittle, Oxid. Met. 18 (1982) 245–284.

19] Y.R. He, D.L. Douglass, F. Gesmundo, Oxid. Met. 37 (1992) 413–439.20] D.J. Young, W.W. Smeltzer, J.S. Kirkaldy, Metal Trans. 6A (1975)

1205–1215.