Nuclear Engineering and Design 206 (2001) 337–350

Copper precipitates in 15 NiCuMoNb 5 (WB 36) steel:material properties and microstructure, atomistic simulation,

and micromagnetic NDE techniques

I. Altpeter a,*, G. Dobmann a, K.-H. Katerbau b, M. Schick b, P. Binkele b,P. Kizler b, S. Schmauder b

a Fraunhofer-Institut fur Zerstorungsfreie Pruf6erfahren IZFP, Uni6ersitat, Geb. 37, 66123 Saarbrucken, Germanyb MPA Stuttgart, Germany

Accepted 24 November 2000

Abstract

The material investigations presented confirm the results of earlier MPA investigations that the service-inducedhardening and decrease in toughness in WB 36 materials are caused by the precipitation of copper. In the initial stateof the material, generally only a part of the alloyed copper is precipitated. The other part is still in solution and canbe precipitated during long-term operation at temperatures above 320–350°C. The copper precipitation leads to adistortion of the crystal lattice surrounding the copper precipitates and yields internal micro-stresses. If the numberand size of the copper precipitates change during operation of a component, a change of the residual-stress leveloccurs. Formation and growth of copper precipitates was simulated using atomistic calculations. In addition, it waspossible to mathematically follow the movement of dislocations and their attachment to precipitates. In this way thenano-simulation was established as a scientific method for the numerically based understanding of precipitationhardening. The results obtained from load stress-related Barkhausen noise measurements demonstrated that thesemicro-magnetic procedures are generally suitable for the verification of copper precipitation. The goal of currentresearch is to establish these findings statistically through further experimental measurements. In addition, theinfluence of different deformation states, macro residual stress, and thermal-induced residual stress have to beresearched. This is important for future developments of non-destructive inspection techniques applied to inservicecomponents. © 2001 Elsevier Science B.V. All rights reserved.

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1. Introduction

The low-alloy, heat-resistant steel 15 NiCu-MoNb 5 (WB 36, material number 1.6368) is used

as piping and vessel material in boiling waterreactor (BWR) and pressurized water reactor(PWR) nuclear power plants in Germany. Onereason for its wide application is the improved0.2% yield strength at elevated temperatures. Inaddition, heat treatment is economical because aferrite–bainite structure with relatively high

* Corresponding author. Tel.: +49-681-93023827; fax: +49-681-93025920.

0029-5493/01/$ - see front matter © 2001 Elsevier Science B.V. All rights reserved.

PII: S0 029 -5493 (00 )00420 -9

I. Altpeter et al. / Nuclear Engineering and Design 206 (2001) 337–350338

bainite content and without any pearlite also re-sults from air cooling after austenitizing. This ispossible due to the nickel and molybdenum con-tents of the steel.

Conventional power plants use this material atoperating temperatures of up to 450°C, whereasGerman nuclear power plants use the materialmainly for pipelines at operating temperaturesbelow 300°C and in some rare cases in pressurevessels up to 340°C (e.g. a pressurizer in aPWR). Following long hours of operation(90 000–160 000) some damage was seen in pip-ing systems and in one pressure vessel of con-ventional power plants during 1987–1992(Adamsky et al., 1991, 1996; Rau et al., 1991;Jansky et al., 1993), which occurred during oper-ation and in one case during in service hydro-testing. In all damage situations, the operatingtemperature was between 320 and 350°C. Eventhough different factors played a role in causingthe damage, an operation-induced hardening as-sociated with a decrease in toughness was seenin all cases. The latter is mainly a shift in thetransition temperature of the notched-bar impacttest to higher temperatures. According toAdamsky et al. (1996), the processes that lead tothe shift in the transition temperature of thenotched-bar impact test in WB 36 are unknown.

Since the beginning of the 1990s, the Materi-als Testing Institute (MPA) Stuttgart has per-formed several projects to shed light upon theoperation-induced hardening and decrease intoughness of WB 36 (Ruoff and Katerbau, 1992;Schick and Wiedemann, 1995; Willer and Kater-bau, 1995; Schick, 1996; Schick et al., 1998).This resulted in obvious indications that theseprocesses are caused by the copper precipitationthat occurs during long-term operation at tem-peratures from 320°C upwards. Based on thislevel of information, a government-supportedproject was started (BMWi-Vorhaben, 1997).The objectives of this project are to describequantitatively the alterations of the mechanicalproperties of WB 36 which are possible underlight water reactor conditions, and to under-stand the underlying microstructural processes.

In parallel to the above activities, MPA Stutt-gart executed calculations using a mesoscopic

theory of precipitation hardening and electron-microscope data, which showed a quantitativecorrelation between the number and size of theservice-related generation of copper precipitatesand the service-related increase of the elasticlimit of WB 36 (Kizler et al., 1998, 2000;Uhlmann et al., 1998). The current project in-cludes an in-depth, theoretically improved, un-derstanding of the mechanical behavior of WB36. The goal is to determine the mechanicalproperties of the material starting from itsatomic structure, i.e. starting from model calcu-lations using the atomic length scale (‘nano-sim-ulation’). The following points are particularlyimportant:� Simulation of the formation and growth of

copper precipitates in steel;� Simulation of the movement of dislocations

and their pinning at precipitates.At IZFP Saarbrucken, projects have been and

are being pursued with the goal to detect thegeneration of service-induced copper precipita-tion in WB 36 using micro-magnetic test proce-dures (Altpeter et al., 1998; BMWi-Vorhaben,1998). A process is being used for the determi-nation of the high degree of higher-order inter-nal stresses that was developed and patentedwithin the framework of a DFG-research projectfor steel containing cementite (Altpeter et al.,1995). The feasibility study (Altpeter et al.,1998) showed that this process could generallybe extended to the determination of high degreeinternal stress for copper precipitates. This pro-ject (BMWi-Vorhaben, 1998) included the statis-tical determination of the usability of theprocedure to verify copper precipitates and toprovide requirements for non-destructive deter-mination and (possible) quantification of WB 36material alterations such as service-induced hard-ening and decrease in toughness.

This paper outlines the current understandingof copper precipitation in WB 36. The three top-ics mentioned above will be included, as follows:� Material properties and microstructure;� Results of the atomistic simulation;� Non-destructive testing using micro-magnetic

techniques.

I. Altpeter et al. / Nuclear Engineering and Design 206 (2001) 337–350 339

2. Material properties and microstructure

Tables 1 and 2 show the specified data (status1965) for the chemical composition and tensiletests of the WB 36 material in comparison to twoother typical steel types used for vessel construc-tion, 15 Mo 3 and 13 CrMo 44. The copper,nickel, and niobium contents, as well as the in-creased values of manganese along with decreasedchromium content are typical for WB 36 material.In the meantime, also the contents of aluminum(0.015–0.050%) and nitrogen (max. 0.02%) werespecified. There are significant differences withrespect to the strength values: the yield strength at350°C for WB 36 is almost twice the value of theother two steels. This, in addition to the econom-ical heat treatment, explains the extensive applica-tion of the WB 36 material in the mediumtemperature range.

Table 3 shows data on the change of materialproperties of WB 36 material components afterlong-term service temperatures ranging from 330to 350°C. The first three components listed inTable 3 have shown damages (Adamsky et al.,1991, 1996; Rau et al., 1991; Jansky et al., 1993).The strength parameters increased up to 140 MPaduring operation, with a corresponding increasein hardness, and the transition temperature forthe notched-bar impact energy was elevated up to70 K. Fig. 1 shows the respective notched-barimpact energy versus temperature curve (KV–Tcurve) for the initial state and after service at350°C for 57 000 h. In addition to the shift of thetransition temperature a reduction of the uppershelf energy, although less relevant, of approxi-mately 20% can be observed. Similar changes are

found for the other components shown in Table3.

It is interesting that the KV–T curve of theinitial state can, for the most part, be restored byrecovery annealing at 550°C for 3 h of the mate-rial state after long-term service, as shown in Fig.1. Further investigations have demonstrated thatthe difference between the initial state and therecovery annealed condition, which is still visiblehere, is virtually lost when recovery annealing isperformed at the temperature of the last stress-re-lief heat treatment of the component (in thepresent case at 580°C).

The dependence of the stress intensity factor KIJ

for crack initiation on the test temperature isshown in Fig. 2. A comparison is made betweenthe initial state and the material condition afterservice (350°C, 57 000 h). The service exposureresults in a shift of the transition temperature ofapproximately 100 K, which is significantly largerthan the shift of 58 K measured by the notched-bar impact tests in this case. The upper shelf valuealso shows a noticeable reduction of the values ofKIJ. Further investigations are being made in or-der to verify these results.

The explanation of the service-induced changesof the properties of the WB 36 material can bederived from the currently available knowledge onthe iron–copper-phase diagram depicted in Fig. 3.The solubility of copper in steel at temperaturesbelow approximately 650°C was unknown untilthe 1980s (Hansen and Anderko, 1958;Kubaschewski, 1982), and it was assumed, obvi-ously, that steel heat-treated at temperatures be-tween 650 and 550°C would not contain anydissolved copper (Haarmann and Kalwa, 1986). It

Table 1Chemical composition of vessel steels (status 1965)a

C Si Mn Cr Mo Ni NbMaterial Cu

Min. 0.12 0.15 0.5015 Mo 3 (TH 31) 0.250.700.350.20Max. 0.35

Min. 0.10 0.15 0.4013 CrMo 4 4 (TH 32) 0.70 0.40Max. 0.18 0.35 0.70 1.00 0.50

15 NiCuMoNb 5 (WB 36) Min. 0.500.25 0.80 0.25 1.000.40 1.30 :0.20 0.800.17 0.50 1.20 0.30Max.

a All values in mass %.

I. Altpeter et al. / Nuclear Engineering and Design 206 (2001) 337–350340

Table 2Mechanical properties of vessel steels (status 1965)

Material Yield strength, ReH Ultimate tensile strength, Rm (RT) (MPa) Uniform elongation, A5 (%)(MPa) at

350°CRT

]17515 Mo 3 (TH 31) 430–520]265 ]19]295 ]215 430–55013 CrMo 4 4 (TH ]18

32)]430 ]353 610–76015 NiCuMoNb 5 ]16

(WB 36)

is now known, however, that WB 36 materialsannealed in this temperature range still containnoticeable amounts of copper in solid solution(Sundman et al., 1985). Therefore, the desiredincrease in strength of the WB 36 material iscaused by only a part of the alloyed copper. Theother part of the copper, which is still in solidsolution in the initial state of the material, slowlyprecipitates during long-term operation at temper-atures above 320°C and may lead to an undesir-able increase in hardness and decrease intoughness. These precipitation processes havelong been described, by the way, but not forsteels, but rather for iron–copper model alloysonly (Hornbogen and Glenn, 1960; Hornbogen etal., 1966; Goodman et al., 1973; Othen et al.,1991, 1994).

First results of service-related copper precipita-tion in WB 36 materials were outlined by Ruoffand Katerbau (1992). Further investigations werecarried out by Willer and Katerbau (1995) andSchick et al. (1998). Fig. 4 shows an example of atransmission electron-microscope (TEM) image ofservice-exposed WB 36 material. Copper particlesranging from 2 to 20 nm in size are present. Someof the particles have a twin structure recognizableby a typical striping. With the TEM investiga-tions, it was furthermore shown that the copperparticles have three differing crystal structuresdepending on their size: up to 6 nm, body-cen-tered cubic, the same as the surrounding matrix,above 20 nm, face-centered cubic like pure cop-per, and between about 6 and 20 nm size, atransition structure. Particles of the latter groupmake up about 50% of all particles visible by

TEM and lead to a pronounced distortion of thecrystal matrix in the regions around the particles,visible under suitable TEM conditions. Therefore,well defined internal micro-stresses are present.

Fig. 5 shows the number and size distributionof the precipitated copper particles visible in theTEM. The initial state of a WB 36 material isshown on top, the material after service at thebottom. Approximately the same volume of mate-rial was analyzed in both cases. The location ofthe maximum of the size distribution and theaverage diameter do not differ significantly. Incontrast, the total number of particles after ser-vice is approximately twice as high as in the initialstate. The results of the TEM investigations showthat the number of particles increases during ser-vice for all of the three size groups. In comple-

Fig. 1. WB 36 — absorbed energy versus temperature curve.

I. Altpeter et al. / Nuclear Engineering and Design 206 (2001) 337–350 341

Tab

le3

WB

36-a

lter

atio

nof

mat

eria

lpr

oper

ties

duri

ngse

rvic

e

Ult

imat

ete

nsile

stre

ngth

,D

Rm

Pow

erpl

ant

Com

pone

ntV

icke

rsY

ield

stre

ngth

,D

Rp

0,2

Serv

ice

Tra

nsit

ion

tem

pera

ture

,D

T(M

Pa)

(K)

cond

itio

nsha

rdne

ss,D

HV

(MP

a)

T-C

ompo

nent

55L

igni

te(1

987)

330°

C/9

160

0h

4884

–82

2840

330°

C/1

6300

0h

Boi

ler

drum

110

Oil-

fired

(199

0)–

50–6

0–

338°

C/1

2800

0h

–P

ipin

gL

igni

te(G

reec

e,19

92)

NP

P(1

993)

–45

339°

C/1

6000

0h

110

Pre

ssur

izer

110

125

5570

Coa

l-fir

ed(1

995)

Boi

ler

drum

350°

C/5

700

0h

140

I. Altpeter et al. / Nuclear Engineering and Design 206 (2001) 337–350342

Fig. 2. WB 36 — fracture toughness at crack initiation.

Fig. 4. WB 36 — precipitations after long-term service.

distribution of the copper precipitates are virtu-ally the same for a material in the initial state andthe material after service and an additional recov-ery annealing (at the temperature of the laststress-relief heat treatment).

mentary investigations using SANS (Small AngleNeutron Scattering) technique it was additionallyfound that the number of particles (not detectablein TEM) having a size of approximately 1 nm alsoincreases significantly. Furthermore, it wasdemonstrated by TEM that the number and size

Fig. 5. WB 36 — size distribution of the copper precipitates.Fig. 3. Iron–copper phase diagram.

I. Altpeter et al. / Nuclear Engineering and Design 206 (2001) 337–350 343

Fig. 6. Results of a Monte Carlo simulation on formation and growth of copper precipitates; only the copper, not the iron atomsare shown.

3. Atomistic simulation (nanosimulation) of WB36 material

3.1. Simulation of formation and growth ofcopper precipitates in steel

The objective of these investigations was todescribe the formation and growth of copper pre-cipitates on the atomic level. Results concerningthe thermal-induced hardening of copper alloyedsteel and the possibility to pre-calculate this quan-tity were expected. Calculations were done usingthe Monte-Carlo simulation program (Soisson etal., 1996), which was modified for the currentinvestigation.

The simulation uses a body centered cubic crys-tal lattice. The model volume used was a cubehaving an edge length of 32 lattice constants andperiodic boundary conditions. Therefore, the lat-tice contains 2×323=65 536 lattice points. In theoriginal state the lattice has 0.6 at.%=393 ran-domly distributed copper atoms, a vacancy and65 142 iron atoms. The simulation has a vacancyconcentration of CSIM=1.526×10−5 but in real-ity, at 350°C, the concentration is CREAL=1.14×10−13. Therefore, the time scale of the simulationis multiplied by the appropriate factor.

The diffusion of the atoms proceeds via va-cancy jumps towards nearest neighbor atoms.This type of position exchange is thermally acti-vated and the jump frequency G is given by:

Gi6=ni exp�−DEi6

kT�

with k the Boltzmann’s constant and T,temperature.

The attempt frequencies (ni) to change the lat-tice location is dependent on the lattice constant aand the diffusion coefficient DO,i of iron or coppervia ni=DO,i/a2. The activation energy DEin is thedifference in energy between the stable state andthe saddle point position of a diffusing atom,which is located next to a vacancy. The activationenergy depends on the local short-range-order andis determined separately for each positionexchange.

The following illustration shows the results of asimulation at a temperature of T=350°C (=623K) as displayed in Fig. 6. After t=2.2 years,small precipitates (approximately 1 nm in diame-ter) have formed from the originally randomlydistributed copper atoms. After t=12.1 yearsthese have merged into four larger precipitates(approximately 2 nm in diameter).

3.2. Simulation of the mo6ement of dislocationsand their pinning at precipitates

The pinning of dislocations at obstacles is thebasic mechanism of precipitation strengthening(Nembach, 1997). In the case of precipitates thatcan be cut, for example copper precipitates in WB36 material, the movement of dislocations is notcompletely blocked, just hindered (Russell andBrown, 1972). During cutting of the precipitate bythe moving dislocation, a characteristic angle isfound between the dislocation segments at themoment of maximum stress, called the critical

I. Altpeter et al. / Nuclear Engineering and Design 206 (2001) 337–350344

Fig. 7. Cross-section through the iron structure model used for the nano-simulation of the dislocation movement; scale in 1 A, (0.1nm).

angle. This is a key parameter for the understand-ing of precipitation hardening. If it can be deter-mined using the nano-simulation, it then providesthe possibility to understand and pre-calculate theprecipitation hardening starting from the atom-istic properties (Russell and Brown, 1972; Nem-bach, 1997).

The model for the present simulation was acuboid-shaped iron mono-crystal, consisting of82 600 atoms, measuring 7.5×7.6×17.0 nm3. Across-section is shown in Fig. 7. Further detailsare published separately (Nedelcu et al., 1999a,b,2000). Two edge dislocations were introduced intothe model, located at the right and left of Fig. 7.Gliding plane and Burgers vector (b) have thecrystallographic label (1–10) and b=9�1 1 1�.

Iron atoms were replaced with copper atoms ina spherical area representing a coherent precipi-tate. In the present case, the copper precipitatehas a diameter of 3.0 nm and consists of 1254copper atoms. Fig. 7 shows a cross-sectionthrough the center of the model; the copper pre-cipitate is indicated in light gray. The model hasinternal stress due to the dislocation and it mustrelieve this stress using a relaxation algorithm. Itturned out that the internal stress was sufficient toinitiate movement of the dislocations. This modelwas retained since the current simulation calcula-tions are the first of this kind. In addition to themovement of the dislocations, the lattice alsorelaxes slightly around the precipitate to accountfor the differing atomic diameters of iron andcopper. This stress is negligible compared to thestress caused by the dislocation.

The molecular dynamics program FEAt wasused for the relaxation of the model (Kohlhoffand Schmauder, 1989). The inter-atomic forces

are represented by the current Embedded-Atom-Model, EAM. The EAM potentials for iron andcopper can be found in the literature. A potentialdeveloped by MPA Stuttgart was used for theiron–copper (Fe–Cu) interaction (Ludwig et al.,1998). Using this relaxation algorithm, the move-ment of the dislocation is initiated without exter-nal force. To investigate this, an algorithm wasdeveloped that recognized the changing positionof the dislocation in the relaxing model, by calcu-lating the Burgers vector density distribution(Nedelcu et al., 1999a,b, 2000). The positions ofthe dislocation lines represent the maxima of thedistribution.

Fig. 8. Dislocation movement: movement to a precipitate (greycircle) and cutting of the precipitate.

I. Altpeter et al. / Nuclear Engineering and Design 206 (2001) 337–350 345

Fig. 8 shows six different stages of the disloca-tion movement relative to a precipitate. The dislo-cations are perpendicular to the cross-section ofthe simulation model depicted in Fig. 7. The firstpartial illustration shows the dislocation at a largedistance from the precipitate (gray circle). Thesecond illustration shows the protuberance of thedislocation line indicating the attraction betweenthe dislocation and the precipitate. The last illus-tration shows the state of maximum stress. Theangle between the dislocation segments, the criti-cal angle, is approximately 140°. This value corre-sponds to the value that can be calculated forcopper precipitates of this size and is confirmedby continuum mechanics theory (Russell andBrown, 1972; Nedelcu et al., 1999a,b, 2000).While various theoretical approximations andparameters (that must be developed through ex-periments) are required for continuum mechanicstheory, the current calculations are based onnothing more than the elastic properties on theatomic length scale as represented by the EAMpotentials. The hardening of the investigatediron–copper system is therefore based on themodulus of transverse elasticity of matrix andprecipitate. The nano-simulation could thereforebe established as a scientific method for the nu-merically based understanding of precipitationhardening using the performed calculations.

4. NDT using micro-magnetic techniques

4.1. Principles

The micro-magnetic concept used is based onload–stress dependent Barkhausen noise measure-ments. The maximum of the Barkhausen noiseamplitude is recorded as a function of the increas-ing load stress. This curve runs through a magne-tostrictive-related maximum. The shift of thismaximum along the stress axis is a measurementof the change of the micro (or macro) residualstress condition. This technique permits the collec-tion of quantitative data of residual stress varia-tions without the use of a reference method suchas X-ray diffraction.

4.2. Experimental set-up

4.2.1. Recording of hysteresis cur6esCylindrical samples are inserted into the Hys-

trometer (manufactured by List) for magnetiza-tion. The magnetizing frequency ( fE) is 1 Hz andthe energized field amplitude (HMAX) is 50 Acm−1. These magnetized samples are then usedfor the recording of ferro-magnetic hysteresiscurves. The voltage induced in an encircling coil isused to calculate the respective induction B(t) byintegration. Previous investigations, jointly per-formed with PTB in Braunschweig, have demon-strated the comparability of the magnetizingcurves using the apparatus with a tolerance ofB10% of PTP’s calibration standard, which issufficient for these materials.

4.2.2. Integral load-stress related Barkhausennoise measurements

To establish load–stress related Barkhausennoise measurements, a measurement system wasinstalled to record micromagnetic test values dur-ing concurrent tensile loading. The samples aremagnetized in the longitudinal direction. TheBarkhausen noise signal is recorded using twodifferential air-core coils to separate the influenceof the energizing magnetic field. The magneticBarkhausen noise is triggered by an alternatingmagnetic field applied to the sample using anelectromagnet. The noise signal is recorded asinducted voltage, appropriately filtered, rectified,amplified, and displayed as function of the tan-gential field strength. The resulting so calledBarkhausen noise profile is evaluated with respectto their maxima and their respective magneticfield positions by a computer (Altpeter et al.,1998).

4.2.3. Recording of the longitudinalmagnetostriction cur6es

All dimensional changes in the ferromagneticmaterial that result from changes in the magne-tization state are called magnetostriction. A dif-ferentiation is made between the longitudinalmagnetostriction (lL) and the transverse magne-tostriction (lT), length variations parallel or per-pendicular to the field direction, and volumetric

I. Altpeter et al. / Nuclear Engineering and Design 206 (2001) 337–350346

Fig. 9. Ferromagnetic hysteresis curve for the service-exposedmaterial state, measured without load stress. The materialstate after recovery annealing showing the same hysteresiscurve.

The hysteresis curves of the service-exposed andrecovery-annealed states are identical (Fig. 9). Adifference of 140 MPa change in yield strength isnot reflected in the B–H-loop. The coercivestrength (Hc) is 6.590.1 A cm−1 in both mi-crostructure states (service-exposed material stateand recovery annealed state). The strong localizedresidual stress of higher order, due to copperprecipitation, is not reflected in the macroscopichysteresis curve since this is averaged overvolume.

The magnetic Barkhausen noise profile and thehysteresis curves were recorded using the encir-cling coil technique (Fig. 10). Both material con-ditions can be differentiated in the unstressedcondition using the maximum Barkhausen noiseamplitude MMAX. This can be explained by thefact that the tension stress sensitivity (load andinternal stress) of the test value MMAX (Altpeter etal., 1995) is much larger than the coercive fieldstrength of the hysteresis curve. The change in themagnetic Barkhausen noise peak position betweenthe service exposed and recovery annealed sampleis within the error tolerance.

The longitudinal magnetostriction behavior ofthe two different material conditions was investi-gated in addition to the Barkhausen noise profile.The dependence of the longitudinal magnetostric-tion lL on the tangential component Ht of theexciting magnetic field is shown in Fig. 11 withoutadditional load. An obvious difference can be

magnetostriction. The value recorded in the areaof magnetic saturation is called saturation magne-tostriction (lS). This is a combination of longitu-dinal and transverse magnetostriction (Bozorth,1951).

A positive magnetostrictive material, like iron-based materials, under the influence of externalstresses shows the following behavior: tensilestress results in the alignment of the magnetizingvectors of the individual domains in the directionof tension (Bozorth, 1951).

The measurement of the longitudinal magne-tostriction under load stress is performed in astandard way using a strain gauge affixed to thesamples and an amplifier (manufactured byHBM). The load stress applied to the sample ismeasured using a second amplifier connected to aload cell mounted on the tensile test machine.

4.3. In6estigation results

Measurements were taken on two materialstates of WB 36 from the same melt, the service-exposed state (‘B’) and the recovery annealedstate (‘E’). The investigated samples consisted ofrods with 8 mm diameter and 160 mm length.TEM investigations have shown that in both con-ditions, three groups of copper precipitates arepresent whereby state B has more copper precipi-tates than state E in every group.

Fig. 10. Barkhausen noise profile curves for the service-ex-posed material state and after recovery annealing, measuredwithout load stress.

I. Altpeter et al. / Nuclear Engineering and Design 206 (2001) 337–350 347

Fig. 11. Variation of longitudinal magnetostriction lL withmagnetic field Ht for the service-exposed material state andafter recovery annealing, measured without load stress.

Fig. 12 shows that the longitudinal magne-tostrictive curve shifts to smaller values underincreasing tensile load, since the density of move-ments at the [100]-90° domain wall decreases inthe direction of magnetization. With increasingtensile loads, the magnetic reversal process trans-forms into magnetostrictive negative rotation (Bo-zorth, 1951). Above the tensile stress limit, thetotal longitudinal magnetostrictive curve is in thenegative area, a characteristic for every type ofsteel. This stress corresponds to the value on theMMAX(s) curve where the maximum occurs. Thisis the case for the service-exposed material statesstarting at sB:55 MPa (top right in Fig. 12) forthe recovery annealed material state, but startingat sE:70 MPa, bottom right of Fig. 12, wherethe stress difference is also Ds:20 MPa. Theevaluation of the longitudinal magnetostrictioncurves therefore yield qualitatively and quantita-tively the same result as the magnetic Barkhausennoise profile curves.

5. Conclusion

The material investigations presented confirmthe results of earlier MPA investigations that theservice-induced hardening and decrease in tough-ness in WB 36 materials are caused by the precip-itation of copper. In the initial state of thematerial, generally only a part of the alloyedcopper is precipitated. The other part is still insolution and can be precipitated during long-termoperation at temperatures above 320°C.

The copper precipitation leads to a distortion ofthe crystal lattice surrounding the copper precipi-tates and yields internal micro-stresses. If thenumber and size of the copper precipitates changeduring operation of a component, a change of theresidual-stress level occurs.

Formation and growth of copper precipitateswas simulated using atomistic calculations. In ad-dition, it was possible to mathematically followthe movement of dislocations and their attach-ment to precipitates. In this way the nano-simula-tion was established as a scientific method for thenumerically based understanding of precipitationhardening.

observed in the two material conditions, lLMAX,E

(4.2490.06 mm m−1) and is much larger thanlLMAX,B (3.4890.05 mm m−1), where lLMAX isthe maximum of the longitudinalmagnetostriction.

The copper-dependent residual stress reducedby recovery annealing leads to a change in the lL

(Ht) curve for the no-load stress condition. Toquantify the degradation of the residual stress of ahigher order, load stress-related micro-magnetictests were carried out.

Both material conditions ‘B’ and ‘E’ can beclearly differentiated in the load stress dependencyof the magnetic Barkhausen noise profile, asshown in Fig. 12. The maximum of the MMAX(s)curve is shifted by approximately Ds:20 MPa,from sB:50 MPa to sE:67 MPa. The curveswere approximated using polynomials of the 6thorder and the maxima of the respective approxi-mation curves were determined analytically. Themagnetic Barkhausen noise profile is influencedby the total stress state of the material, i.e. thesuperposition of macro stress and stress of ahigher degree. If internal stresses, due to copperprecipitates, are reduced by recovery annealing,then a higher load must be applied to arrive at thesame total stress condition. The stress differenceof Ds:20 MPa indicates an integral value for thetotal sample volume. Local stress concentrationscan be much higher at the phase boundaries ofcopper and matrix.

I. Altpeter et al. / Nuclear Engineering and Design 206 (2001) 337–350348

The results obtained from load stress-relatedBarkhausen noise measurements demonstratedthat these micro-magnetic procedures are gener-ally suitable for the verification of copper precipi-tation. The goal of current research is to establish

these findings statistically through further experi-mental measurements. In addition, the influenceof different deformation states, macro residualstress, and thermal-induced residual stress has tobe researched. This is important for future devel-

Fig. 12. Variation of the maximum MMAX of the Barkhausen noise amplitude and of longitudinal magnetostriction lL with appliedload, for material states after service (upper row) and after recovery annealing (lower row).

I. Altpeter et al. / Nuclear Engineering and Design 206 (2001) 337–350 349

opments of non-destructive inspection techniquesapplied to inservice components.

Acknowledgements

The authors would like to thank the (German)Federal Ministry of Education and Research(BMBF), for Economic Affairs (BMWi) and forEnvironmental Protection, Nature Conservation,and Reactor Safety (BMU) for the promotion andsponsorship of the research work. In addition, wethank Dr F. Soisson and Professor Dr G. Martin(both CEA Saclay) for providing the Monte-Carlo software and their support during the fur-ther development of the software.

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