NMIlSlll'imiii »„ Sr*N° * TECH R.I.c.

IIIDS M7fiflflL|

NtST

f'^BUCATlOA/s

NisrUnited States Department of Commerce

Technology Administration

National Institute of Standards and Technology

rJ/ST Technical Note 7500-4

Materials Reliability Series

oc

100

.115753

NO. 1500-4

1998

Nondestructive Characterization of

Reactor Pressure Vessel Steels:

A Feasibility Study

Final Report

NRC Program RES-97-005

NIST Technical Note 1500-4

Materials Reliability Series

Nondestructive Characterization of

Reactor Pressure Vessel Steels:

A Feasibility Study

Submitted by

Harry I. McHenry, Program Manager

George A. Alers, Principal Investigator

Materials Reliability Division

Materials Science and Engineering Laboratory

National Institute of Standards and Technology

325 Broadway

Boulder, Colorado 80303-3328

Final Report, NRC Program RES-97-005

Control Number: B&R NO. 76015114010

Job Number: W 6738

Appropriation Symbol: 31x0200.760

BOG 253A

June 1998

'^TES O*

U.S. DEPARTMENT OF COMMERCE, William M. Daley, Secretary

TECHNOLOGY ADMINISTRATION, Gary R. Bachula, Acting Under Secretary for Technology

NATIONAL INSTITUTE OF STANDARDS AND TECHNOLOGY, Raymond G. Kammer, Director
National Institute of Standards and Technology Technical Note

Natl. Inst. Stand. Technol., Tech. Note 1500-4, 104 pages (June 1998)

CODEN:NTNOEF

U.S. GOVERNMENT PRINTING OFFICE

WASHINGTON: 1998

For sale by the Superintendent of Documents, U.S. Government Printing Office, Washington, DC 20402-9325
CONTENTS

1. Introduction 1

2. Measurement Results 3

2.1. Materials 3

2.1.1 RPV Alloy A533B 3

2.1.2 Dilute Copper in Iron Alloys 4

2.1.3 Preparation of Materials and Test Specimens 5 2.1.4 References 6

2.2. Mechanical Properties 16

2.3. Elastic Properties 20

2.3.1 Introduction 20

2.3.2 Details of Measurement Methods 20

2.3.3 Results and Discussion 21

2.3.4 Conclusions 23

2.3.5 Recommendations 23

2.3.6 References 24

2.4. X-ray Diffraction 44

2.4.1 Methodology. 44

2.4.2 Experimental Technique 44

2.4.3 Results 44

2.4.4 Physical Model 45

2.5. Nonlinear Elastic Properties 47

2.5.1 Background 47

2.5.2 Experiments 47

2.5.3 Results 47

2.5.4 Discussion 48

2.5.5 References 49

2.6. Magnetic Properties 52

2.6.1 Introduction 52

2.6.2 Literature Review 52

2.6.3 Vibrating Sample Magnetometer (VSM) Measurements 52

2.6.4 Apparatus 52

2.6.5 Samples and Sample Preparation 53

2.6.6 Data and Calibration Runs. 53

2.6.7 Data Analysis. 53
2.6.8 Low Temperature VSM Experiments 54

2.6.9 Results and Conclusions. 55

2.6.10 Other Methods 56

2.7. Micromagnetic Techniques 58

2.7.1 Development of a Noncontacting Measurement of Magnetostrictioa 58

2.7.2 Measurement of A7 10 Steel 61

2.7.3 Measurements on A533B.Steel 61

2.7.4 Conclusions 62

2.7.5 References 63

2.8. Dislocation Mobility Under a Static Stress 72

2.8.1 Introduction 72

2.8.2 Experimental Apparatus 72

2.8.3 Results 73

2.8.4 Conclusions 75

2.8.5 References 75

2.9. Measurements Under Cladding 82

2.9.1 Magnetic Field Penetration of SNUPPS Cladding 82

2.9.2 Static Magnetic Properties (B-H Looper) 84

3. Discussion 88

3.1. Physical Property- Strength Correlations 88

3.2. Dislocation Mechanisms 90

3.3. Temperature Effects 91

4. Conclusions 96

4.1. Task I 96

4.2. Task II 96

4.3. Task III 97

5. Recommendations 98
FOREWORD

The Materials Reliability Series of NIST Technical Notes are reports covering significant

research accomplishments of the Materials Reliability Division. The Division develops

measurement technologies that enable the producers and users of materials to improve the quality

and reliability of their products. Measurement technologies are developed for process control to improve the quality and consistency of materials, for nondestructive evaluation to assure quality

of finished materials and products, and for materials evaluation to assure reliable performance.

Within these broad areas of measurement technology, the Division has focused its resources on

three research themes:

• Intelligent Processing of Materials—To develop on-line sensors for measuring the

materials' characteristics and/or processing conditions needed for real-time process

control.

• Ultrasonic Characterization of Materials—To develop ultrasonic measurements for

characterizing internal geometries of materials, such as defects, microstructures, and

lattice distortions.

• Micrometer-Scale Measurements for Materials Evaluation—To develop measurement

techniques for evaluating the mechanical, thermal, and magnetic behavior of thin films

and coatings at the appropriate size scale.

This report is the fourth in the Materials Reliability Series. It covers research on nondestructive

characterization of RPV steels, which is part of our research on ultrasonic characterization of

materials. Previous reports in this series are:

Technical Note 1500-1 Tensile Testing of Thin Films: Techniques and Results,

byD.T. Read, 1997

Technical Note 1500-2 Procedures for the Electron-Beam Moire Technique,

by E.R. Drexler, 1998

Technical Note 1500-3 High-Energy, Transmission X-ray Diffraction for Monitoring

Turbine-Blade Solidification, by D.W. Fitting, W.P. Dube,

and T.A. Siewert, 1998

NONDESTRUCTIVE CHARACTERIZATION OF REACTOR

PRESSURE VESSEL STEELS: A FEASIBILITY STUDY

Radiation damage to the walls of reactor pressure vessels (RPVs) causes the steel to become more brittle and less able to withstand the thermal stresses of start-up and shutdown procedures. Current methods of monitoring the degree of embrittlement are based on measurements of the ductile-to-brittle transition temperature (DBTT) of surveillance

specimens subjected to severe radiation damage inside the reactor itself In order to improve

on this conservative approach and extend the useful life of vessels that have been in service

for many years, NIST undertook a feasibility study to investigate nondestructive techniques

for inferring the DBTT of the pressure vessel wall itself The approach used was based on

the hypothesis that the changes in microstructure that accompany embrittlement could be

detected by accurate measurements of the physical properties of the steel in the pressure

vessel wall. Nine magnetic and ultrasonic physical properties were measured on samples

of RPV steel and on an A710 steel that could be heat treated to exhibit the same

microstructural changes and embrittlement features as found in radiation damaged steel. The results showed that a nonlinear elastic coefficient, the magnetostriction constant and

a dynamic magnetic permeability were strongly correlated with microscopic internal strains that appear to inhibit dislocation motion and cause embrittlement of the A710 steel. A study

of techniques for making measurements of these three properties on a pressure vessel as well as for establishing their relationship to radiation embrittlement of A533B reactor steel

is currently underway.

Key words: copper rich precipitates; elastic constants; internal friction; internal strain; magnetostriction; microstructure; nonlinear elasticity

1. INTRODUCTION

This program was undertaken in order to assess the feasibility of using physical property

measurements to evaluate the embrittlement of reactor pressure vessel (RPV) steels. It is based on the premise that there are measurable physical properties that reflect the changes in

microstructure that accompany radiation-induced embrittlement. If such properties can be found,

then techniques and standards can be developed that will allow in-situ nondestructive testing of a

pressure vessel to ensure its safe operation.

The approach proposed by NIST for the feasibility study was to accurately measure

physical properties of unirradiated RPV steels as well as surrogate materials whose embrittlement

mechanisms are similar to those thought to be responsibile for radiation induced embrittlement.

The physical properties chosen were restricted to those that could be applied to the inner surface

of a pressure vessel coated with a 7 mm thick stainless steel cladding. Thus, magnetic and

ultrasonic methods that could penetrate the cladding were emphasized although x-ray diffraction

studies were performed in order to establish the mechanisms behind the correlations observed

between embrittlement parameters and physical-property measurements. Three sequential tasks were undertaken with the aim of producing recommendations for development of nondestructive

testing procedures that could infer the ductile-to-brittle transition temperature of a pressure vessel

embrittled by radiation damage. These tasks were as follows:
Task I. Procure samples of A533B RPV steels that have not been irradiated but which

showed a spectrum of ductile-to-brittle transition temperatures (DBTTs) and perform a full set of

ultrasonic and magnetic measurements on them as a function of temperature. The output of this

task is intended to establish baseline values for the physical properties of RPV steels and to expose any relationships they might have with the DBTT.

Task II. Acurately measure ultrasonic and magnetic properties of a surrogate steel as it is embrittled by mechanisms similar to those expected to occur in irradiation embrittled pressure

vessels. The output of this task is intended to establish those properties that are sensitive to embrittlement in radiation damaged RPV type steels and that can be measured nondestructively.

Task III. Investigate the possibilities for measuring the embrittlement related properties

of the irmer surface of a reactor pressure vessel that is clad with stainless steel.

Because this feasibility study demanded the performance of very accurate measurements

by specialists in particular techniques, the body of the report has been organized by the physical

property that was measured rather than by the tasks listed above. However, the Discussion

Section presents the results as they apply to the tasks and the conclusions section presents the

recommendations for the development of nondestructive testing methods and procedures. The

table below lists the physical properties measured, the NIST staff member involved, and the

rationalization for its relationship to mechanical strength parameters.

Table 1.1 Physical-mechanical properties measured in this study.

Property Specialist Information content

Elastic moduli

Moduli and internal

friction

X-ray diffraction

Magnetic hysteresis

Magnetostriction

Incremental permeability

Nonlinear ultrasonics

Stress induced damping

Tensile properties

Charpy energy

Hardness

S. Kim Direction dependence of the velocity of sound

H. Ledbetter Dislocations and precipitation effects

D. Balzar Internal strain distributions on a submicron scale

F. Fickett Domain wall/imperfection interactions

B. Igarashi Domain wall mobility and internal stresses

B. Igarashi Dynamic magnetic response

D. Hurley Interatomic strain associated with precipitation

W. Johnson Force required to move dislocations

D. McColskey Yield and ultimate strength. Elongation

D. Vigliotti DBTT and impact energy absorption

P. Purtscher Rockwell A hardness
2. MEASUREMENT RESULTS

2.1 MATERIALS —P. T. Purtscher

2.1.1 RPV Alloy A533B

Six samples of RPV steel were sent to NIST from ORNL. Four were cut from the

SNUPPS vessel and two came from the Shoreham vessel. They were designated as shown in

Table 2.1.1.

Figure 2.1 .1 is a sketch that shows how the various test specimens were cut from the two

largest segments of the SNUPPS material. Note that 20 Charpy specimens were cut from the

weld and 20 more were cut from each of the four adjoining plates. These were designated W-1

(from the weld) and 1-1 from Segment 1, Plate 1; 1-2 from Segment 1, Plate 2: etc.

Half of the cladding from the 157 x 66 mm face of SNUPPS Seg. #4 was cut away from

the base metal by EDM techniques to produce a 76 x 57 mm plate sample of the stainless steel approximately 3 mm thick so that its acoustic and electrical properties could be measured

independently of the A533B substrate.

The two full thickness samples cut from the Shoreham RPV were delivered to NIST in

order to provide examples of material coated with the stainless steel cladding and having the

geometry that will be faced in the ultimate application of any nondestructive test. They each

weigh approximately 1 000 kg (2200 lb) and have been mounted on special heavy duty carts in

order to facilitate their being moved to different laboratories within the NIST facilities.

Table 2.1.1 Designations assigned to the RPV materials provided to NIST by ORNL.

Designation Dimensions (mm) Description

SNUPPS Seg. #1

SNUPPS Seg. #2

SNUPPS Seg. #3

SNUPPS Seg. #4

Shoreham I

Shoreham II

775 X 264 X 95 Girth weld and rolling dir. parrallel to 264 dim.

914 X 240 X 50-90 Axial weld parrallel to 240,

rolling dir. along 914 dim.

226 X 157 X 66 Cladding over the 157 x 66 mm face

226 X 175 X 66 Cladding over the 175 x 66 mm face

953 X 914 X 159 6 mm thick cladding over 940 x 914 face

921 X 890 X 159 6 mm thick cladding over 914 x 890 face
2.1.2 Dilute Copper in Iron Alloys

A recent review article [2.1.1] and several previous publications [2.1.2 - 2.1.4] ascribe

the decrease in ductility and embrittlement of the pressure vessel to the formation of copper-rich

precipitates (CRPs) which are known to increase the flow stress and hardness. The CRPs are

typically 1 to 2 nm in diameter and are a precursor to the formation of the equilibrium e phase

shown in the iron-copper phase diagram. The Cu content of pressure vessel steels is normally

never over 0.45 mass percent, but given the typical neutron fluence of 1 to 3 x 10'^ cm^ (for

neutrons with energies over 1 MeV) and a typical Ni content of 1 .4 mass percent, the shift in the

DBTT compared to the pre-irradiated properties can be as large as 230 °C [2.1 .5]. Due to the

small size of the CRPs, the monitoring of their formation is difficult but is very important to the

safety of the RPV.

2.1.3 Preparation of Materials and Test Specimens

In order to find those physical properties that can sense the formation of CRPs and hence

be candidates for a nondestructive predictor of embrittlement, two steels containing 0.35 (SAE

1508) and 1.13 mass percent copper (ASTM A710) from earlier NIST programs were evaluated

for their suitability as surrogate materials. The compositions are included in Table 2.1 .2. The

effects of direct quenching from hot rolling and subsequent tempering on these two steels was

published by Manganello and Wilson [2. 1 .6]. Figure 2.1 .2 shows the effect of tempering

temperature on the yield strength (YS) for 19 mm thick plates from the two steels after direct

quenching. Steel A (1508) demonstrated a definite peak in the YS after tempering at temperatures above 500 °C; steel C (A710) demonstrated a peak in the YS at 510 °C. The peak

in YS is considered to be due to CRPs. The 19 mm thick plates that were available at NIST were

in the peak-aged condition for each composition and hence include the CRPs in the

microstructure.

Our initial experiments with the two potential surrogate steels were to determine whether

the strengthening due to CRPs could be observed in heat treated conditions as well as in the direct

quench and tempered condition. For this purpose, small samples (10 mm x 10 mm x 55 mm)

were heat treated to various conditions. The standard solution treatment involved heating the

samples to 900 °C and holding for 1 hour. This solution treatment will dissolve all of the Cu in

austenite while retaining a 10 |nm austenite grain size. After solution treatment, the samples were

then cooled in vermiculite to room temperature and aged at temperatures between 450 and 700

°C for 1 hour. Hardness was used for the evaluation in this initial experiment. The hardness of

steel A was R^ 50 and did not increase with subsequent aging treatments. For steel C, a definite

peak in hardness was observed at about 525 °C and the minimum hardness was found at 700 °C.

(See Fig. 2.1 .3.) Thus, we concluded that steel C could be heat treated to produce precipitation

strengthening from the CRPs, whereas steel A could not.

For the nondestructive characterization of microstructure, larger samples (19 x 25 x 300

mm long) were prepared so that multiple samples could be machined from each bar. The goal
Table 2.1.2 Nominal composition (in mass percent) of the steels investigated in this study.

Designation Cu Ni Mn Si Cr Mo Nb Al

A(1508) -0.35 0.35

C(A710) 1.13 0.87

A533B 0.11 0.8

0.08 1.35 0.25 0.03 0.06 0.02 0.022

0.05 0.53 0.30 0.75 0.21 0.32 0.023

0.19 1.29 0.2 0.12 0.53 0.015

was to produce underaged, peak-aged, and overaged samples with respect to the precipitation

strengthening due to CRPs for each steel. In steel A, precipitation strengthening could be

achieved only with the as-received, directly quenched and tempered plates. Therefore, the as- received plate, designated A5, would be the peak aged condition and aging of the as-received

plate at 675 °C would represent the overaged condition, designated A6. Underaged, directly

quenched samples were not available so samples quenched in water, designated AA, and air

cooled, designated AA-2, from the 900 °C solution heat treatment temperature were produced.

Both AA samples were underaged with respect to CRP strengthening, but with different

microstructures. For steel C, the large bars were air cooled to room temperature, designated CA

(underaged), and then some bars were aged at 525 °C, designated C5 (peak aged), and others were aged at 700 °C, designated C7 (overaged). A summary of the microstructures evaluated by

the nondestructive characterization methods is shown in Table 2.1.3.

Table 2.1.3 Microstructures observed in the copper alloys that were

investigated following various heat treatments.

DesignationHeat

treatment

Matrix

metallography

State of

copper

Hardness

Rockwell A

AA As quenched Acicular a + Bainite Underaged 59.0

AA-2 Air cooled Polygonal a + Bainite Underaged 50.0 ±0.5

A5 As received Tempered a + 2nd phase Peak aged 54.3 ±0.5

A6 Rec'd + 675 °C* Tempered a + 2nd phase Overaged 53.1 ±0.5

CA Air cool Polygonal a + 2nd phase Underaged 55.6 ±0.5

C5 Air Cool + 525 °C* Polygonal oc + 2nd phase Peak aged 58.5 ±0.5

C7 Air Cool + 700 °C* Polygonal a + 2nd phase Overaged 53.2 ±0.5

*for 1 hour
For steel A (1508), the microstructures are represented by the following micrographs.

Figure 2.1.4 shows the acicular a and bainitic microstructure of sample AA. The water quenching

was not fast enough to produce a martensitic matrix. Figure 2.1.5 shows the polygonal a plus

bainite structure found in the air cooled AA-2 sample. These two represent the conditions with

the least precipitation strengthening (underaged with 0.35% Cu); AA has the highest hardness

because water quenching did produce transformation strengthening. Figure 2.1 .6 is representative

of both A5 and A6, which had very similar microstructures when observed in the light

microscope. This microstructure is mainly tempered a with some second phase scattered

throughout, indicative of the low C content.

For steel C (A710), the microstructures observed in the light microscope for all three

samples are very similar. Figure 2.1.7 shows the air-cooled CA sample, which was not aged

(underaged with 1.13% Cu). The microstructure is mainly polygonal a but with a dispersed

second phase that is probably a granular bainite, described in detail by Thompson and Krauss

[2. 1 .3]. Figure 2. 1.8 shows the microstructure found in C5 (peak aged), which is similar to that

shown in Figure 2.1.7. The formation of CRPs is not visible at this magnification. The presence

of CRPs is inferred from the hardness change that occurs during aging. The fact that the

hardnesses of CA, C5, and C7 do not correspond exactly with those in Figure 2.2.2 reflects the

influence of specimen size (cooling rate) on the matrix strength due to differences in

transformation strengthening.

The microstructure of the A533B is quite different from that observed in both steels A and

C. Figure 2. 1 .9 shows a low magnification view from a representative area in SNUPPS Seg. #1

(plate 1-1). There is some chemical segregation that produces light and dark etching contrast

with a spacing of about 100 |im. At a higher magnification, the etching contrast is due to the

amount of second phase, light contrast being less second phase. Fig. 2.1.10, and the dark contrast

due to more second phase. Fig. 2. 1.11. In either case, the microstructure of the A533B is finer

than that of the surrogate samples and has a higher volume fraction of second phase due to the

higher C content.

After heat treatment of large samples of steels A and C individual test specimens for each

physical property were machined from the blocks. The specimen size for each test is summarized

in Table 2.1.4.

2.1.4 References

2.1.1 Odette, G.R.; G.E. Lucas, An integrated appoach to evaluating the fracture of irradiated

irradiated nuclear reactor pressure vessels. Nondestructive Eval., 15(3/4): 137; 1996.
Table 2.1.4 Specimen dimensions used for the physical properties

measured in this study.

Physical property Shape Specimen dimensions (in mm)

1Magnetic - He and Ms Cylinder

2. Magnetostriction coef. Block

3

Transverse incremental permeability Block

4. C,j and Q"' and density Block

5. Young's Modulus and Q Sq. Bars

6. Shear and longitudinal wave velocity Block

7. Acoustic nonlinearity parameter, P Block

8. Dislocation breakaway stress Cylinder

9. Tensile strength Cylinder

2 (dia) X 20 long & 3 (dia) x 30 long

9 X 20 X 125

9x20x 125

10 X 11 X 12

4x4 X 19, 25 and 50 long

25x50x19

25 x50x 19

20 (dia) X 160 long

12 (dia) X 75 long

2.1.2 Lahiri, S.K.; Chandra, D.; Schwartz, L.H.; Fine, M.E., Modulus and Mossbauer Studies of

Precipitation in Fe-1.67 At. pet. Cu, Trans. Metall Soc. AIME, 245:1865, Sept. 1969.

2.1.3 Thompson, S.W.; Krauss, G., Copper Precipitation During Continuous Cooling and

Isothermal Aging of A710-Type Steels, Metall. Mater. Trans. A, 27A:1573, June 1996.

2.1.4 Aroztegui, J.J.; Urcola, J.J.; Fuentes, M.; The Influence of Copper Precipitation and

Plastic Deformation Hardening on the Impact Transition Temperature of Rolled Structural

Steels, Metall. Trans. A, 20A:1657, Sept. 1989.

2.1.5 United States Nuclear Regulatory Commission, Regulatory Guide 1 .99: Radiation

Embrittlement of Reactor Vessel Materials, U.S. Government Printing Office,

Washington, D.C. 1988.

2.1.6 Manganello, S.J.; Wilson, A.D., Direct Quenching and Its Effects on High-Strength

Armor Plate, in Low-Carbon Steels for the 90's, R. Asfahani and G. Tither, eds. TMS,

Warrendale, PA, p. 235, (1993).
Segment #1

1/2 T

1/4 T

Plate 1 CVN TensUes

Plate 2

CVN TensUes CVN

Girth Weld

1/4 T

Segment #2

Plate 1 CVN Tensiles

Plate 2

Tensiles CVN —RD—

Outer Wall- V-A-yi Axial Weld

Fig. 2.1.1 Sketch of the SNUPPS Segments #1 and #2 showing the location of pieces used for

various measurement specimens. T = thickness, CVN = Charpy V Notch.
900

800-

a. 700

c 600

Ho^500

400-

300-

as- quenched

steel C(A710)

300 400 500 600 700 800

Temperature, °C

Fig. 2.1.2 Yield strength versus tempering temperature for 19-mm-tliick plates of Cu-containing

steels produced by the direct-quench process [2.1.6].

1—'—I—I—I—I—I—'—I—'—I—'—I—'—I—'—

250 300 350 400 450 500 550 600 650 700 750

Aging temperature, °C

Fig. 2.1.3 Hardness versus aging temperature for steel C (A710) after solution heat treatment at 900 °C for 1 h and cooling in vermiculite.
tr

•ft-'

^

>i^ --^i

•w"^

:^

1^. - .-4^

<tf>r ^2^ Vi^"

Fig. 2.1.5 Light micrograph of the air-cooled steel A sample AA. Bar = 20 fim.

10
k f

'

* ••«.,

%

Fig. 2.1.6 Light micrograph of the direct-quenched and tempered steel A sample A5. This

microstructure is also representative of sample A6. Bar = 20 fim.

11
'^^^ '';^r

-J r- Iter, .-iJs^ ^TlL '-<^

• '• -u-^V;,^- * •• " ^^'7>>-f\^^: ^-!U" '--v i>'.\

.20 Mm . Vl/ ^^^'^V^r

^ i ^ ^4PV'

.^\

(a)

*- p

•

V V

}r*- "^^- K/ -X'

f,(^-ri/'^

''U

.7 Lim

'7 JkL^

• i-^^

(b)

Fig. 2.1.7 Light micrographs of the air-cooled steel C sample CA. Bar = 20 jam in (a) and 7

^im in (b).

12
^-v. .^^i/^ ^^^^^-Ti. i'''5^>Mv_^ i-f^rC-nn^^jt -'^ ^-1 -^

..•tf if #vv i£U- f*!"*? f^

(a)

^'"> \

(b)

Fig. 2.1,8 Light micrographs of the air-cooled and aged steel C sample C5. This

microstructure is also representative of sample C7. Bar = 20 |Lim in (a) and 7 |am in (b).

13
\ 100 ^im

Fig. 2. 1 .9 Light micrograph at low magnification of the quenched and tempered A533B base

metal sample 1-1. Bar= 100 |am.

^^/v^C ^1?^ ^ • /

Fig. 2.1.10 Light micrograph at high magnification of the quenched and tempered A533B base

metal sample 1-1, lighter etching area. Bar = 7 ^im.

14
Fig. 2.1.11 Light micrograph at high magnification of the quenched and tempered A533B base

metal sample 1-1, darker etching area. Bar = 7 |im.

15
2.2 MECHANICAL PROPERTIES —Z). McColskey and D. Vigliotti

Standard tensile test specimens were machined and tested according to ASTM

recommended procedure E-8. The A-533B tensile specimens were machined with a test section

diameter of 12.7 mm (0.5 in) and, because of material thickness limitations, the CA, C5, C7, AA,

A5 and A7 materials were machined with a 6.37 mm (0.25 in) test section diameter. The weld

section of the A533B material was also tested in the smaller diameter because of limited material

availability. In all cases, the test section gage length was four times the specimen diameter.

Tensile tests were conducted in a screw-driven tensile testing machine at an initial strain

rate, in the elastic region of the test, of 0.02 per minute. An extensometer monitored the specimen

test section elongation during the test. The strain sensitivity during the yielding process allowed

the detection of a 0.01% deviation from linearity to be detected. Thus, the proportional limit

could be measured even in cases where a discontinuous yield drop was observed. Table 2.2.1 lists the tensile properties deduced from load-elongation curves for all the materials studied in this

program. Many of the materials showed discontinuous yielding behavior so the upper yield and

the 0.2% offset yield strengths are listed in the table. Yield strength was determined using the

0.2% offset method. Yield point, for materials that exhibited discontinuous yielding, was

determined using the autographic diagram method. Those materials that showed no load drop

but continuously work hardened at stresses above the proportional limit have been listed in the

table as having no upper yield point.

Also included in Table 2.2.1 are values for the Rockwell A hardness and the 41 J (30 ft •

Ibf) ductile-to-brittle transition temperature (DBTT). The latter was deduced from measurements on 1 5 standard notched Charpy bars (from each material) machined and tested according to ASTM recommended procedure E-23. Test temperatures for the Charpy specimens were

selected to establish a full transition temperature curve and ranged between +100 and - 100 °C.

The A series alloys did not show significant changes in hardness in response to the heat

treatments intended to form CRTs so a complete set of physical property and mechanical strength

values was not developed for the A alloy series. However, the C alloy series showed clear

evidence for the strengthening effect of the CRPs. Therefore, a very complete set of data that

included the DBTT was generated for this alloy and they are presented in Table 2.2.1 . Figure

2.2.1 shows the low strain portions of the stress-strain curves obtained on the C series alloy

(A710 steel) in the solution treated, peak aged, and overaged conditions. These curves show the

large increase in flow stress as the solution treated or as-cooled alloy was tempered to the peakaged or peak hardness condition produced by the formation of the CRPs. Upon overaging,

copper precipitates form and a discontinuous yielding response develops at a stress level higher

than the proportional limit of the peak aged condition. Since the hardness is often used as a measure of strength and is correlated with various nondestructively measured physical properties,

it is instructive to plot hardness against the other mechanical properties that can be deduced from

a tensile test. Figure 2.2.2 shows such a graph relating the proportional limit, the 0.2% offset

yield strength and the ultimate tensile strength to the hardness. It is clear that there is a good.

16
Table 2.2.1 Tensile properties of the materials used in this study as deduced from

constant strain rate tensile tests on cylindrical samples

Material Prop. Upper yield 0.2% offset Ultimate Elongation Hardness 41 JDBTT

limit point yield strength tensile str. % Rockwell A °C

MPa MPa MPa MPa

CA 173 415 650 27.8 55.6 -33

C5 391 — 551 718 25.1 58.5 -10

C7 446 471 471 558 29.2 53.2 -60

AA 287 575 821 6.6 59

A5 576 576 560 610 24.8 54

A6 524 524 514 594 25.3 53

A533 1-1 479 491 465 610 38 55.3 -20

A533 1-2 461 468 452 574 43 53.3 -60

A533 2-1 476 482 475 600 40 55.1 -25

A533 2-2 513 525 507 640 35 56.0 -20

A533 W-1 568 589 578 672 24 57.0 -30

linear correlation between the hardness and the ultimate tensile strength but no linear correlation

between hardness and the proportional limit (when irreversible plastic flow is initiated) or the

0.2% offset yield strength.

An even more striking display of the effect of heat treatment and the formation of CRPs in

the A 710 steel alloy is shown in Fig. 2.2.3 where the absorbed energy from the Charpy impact

test is plotted against temperature for the three heat treatment conditions. This graph shows that

both the transition temperature and the upper shelf energy are changed by the aging process and

that overaging makes the transition from ductile to brittle very sharp. For comparison purposes,

the results of Charpy impact tests on the unirradiated A533B material are also plotted in Fig.

2.2.3. For the RPV material, the transition is gradual, at a higher temperature and exhibits a low

upper shelf energy compared to that of the 1.13% Cu alloy. With increasing radiation damage,

A533B shows an even lower upper shelf energy and a less pronounced transition between ductile

and brittle behavior.

17
&

u

600 -r

500--

400--

300-

200 --

100 -

UTS=718MPa

UTS=650 MPa

UTS=558 MPa

0.2 0.4 0.6 Strain (%)

0.8

Fig. 2.2. 1 Low strain portion of the stress strain curves for the Fe-1 . 13% Cu alloy in the three

heat treatment conditions.

1.4 C7

Normalized "Strength" vs Hardness

CA C5

cI

1.3

1.2 +

1.1

1

0.9

o.e

0.7 4

0.6

0.5

0.4 +

0.3

UTS

)BTT

N 0.2% Y.S. s

\

Prop. Lim. ^x^

53 54 55 56 Hardness (RA)

57 58 59

Fig. 2.2.2 Graphical representation of Table 2.2.1 showing the relationship between the hardness

and four strength measures: the proportional limit, the 0.2% yield strength, the ductile-to-brittle

transition temperature and the ultimate tensile strength for the Fe-1. 13%) Cu in the overaged

condition (C7), the solution treated condition (CA) and the peak-aged condition (C5).

18
o^u280- «

240- CA

>;200-

©

S 160-

|l20-

C7.

<

•

(0 < 80- ? /: .'^/y^ A533B Steel

40-

0-

R-d ^'/^ BM1-1

-150 -100 -50 50 100 150

Temperature, °C

Fig. 2.2.3 Charpy energy versus temperature for the Fe-1.13% Cu alloy in the three heat

treatment conditions: CA - solution treated, C5 - peak aged, C7 - overaged. Also shown for

reference are measurements on A533B steel.

19
2.3 ELASTIC PROPERTIES —Sudook Kim

2.3.1 Introduction

We report measurements of elastic-stiffness coefficients and their associated imaginary

parts, the internal frictions, for three body centered cubic (b.c.c.) steels, one in three heat

treatment conditions. We used four dynamic measurement methods and varied the strain

magnitude, the frequency, and the temperature. Our measurements also produced values for the mass density.

The elastic-stiffness coefficient, a fourth-order tensor ^ ^^i, contains two parts, real and

imaginary:

^m = ^,= C^O+iQ-'). (1)

The four-index to two-index transformation represents the usual Voigt contraction method. The

internal-friction term Qj] is small: less than 10"'^ and sometimes much smaller, say 10"^ Both

properties, C^ and Q-], correlate strongly with other physical-mechanical properties. C^ relates

strongly to lattice-vibrational properties such as specific heat; thus it is relatively structure

insensitive. However, Q-] is strongly structure sensitive and is used widely to study material

defects such as dislocations, grain boundaries, vacancies, and so on.

2.3.2 Details of Measurement Methods

The five cases we considered (CA, C5, C7, 1-2, Wl) are described elsewhere in this

report in terms of chemical composition, heat treatment, hardness, plastic-deformation behavior,

and so on.

There are numerous methods for measuring the Cy. The most common—slope of

quasistatic stress-strain curve—is also the least accurate, and yields little reliable Q-] information.

Therefore, we used four dynamic methods:

TP, forced-vibration torsional pendulum (near 1 Hz);

MOR, Marx-oscillator resonance (50 to200 kHz);

ARS, acoustic-resonance spectroscopy (50 kHz to 5 MHz);

PES, pulse-echo superposition (5 to 10 MHz).

We made most measurements by the third method. For these materials, the ARS method gives

0.05% uncertainty in Cy and 10% uncertainty in Q-].

20
2.3.3 Results and Discussion

We made all measurements at ambient temperatures, except for E versus T and Q"' versus T for the CA steel measured between 295 and 200 K. For each of the five steels, we constructed

tables of 48 measured physical-property values, thus 240 values altogether. For all cases, we

obtained by analysis some additional physical-property values: the monocrystal

elastic-stiffness coefficients C°j. Tables 2.3.1 through 2.3.3 and Figures 2.3.1 through 14 show

most of our main results.

Compared with typical commercial b.c.c. steels, the elastic-shear stiffnesses show low to moderate anisotropy with values ranging from 1.000 (isotropic) to 1.085, about 6% of the

maximum-possible anisotropy. (See Table 2.3.3.) With aging, the anisotropy decreased slightly

in the C series steels. Both Wl and 1-2 steels show nearly isotropic elastic stiffnesses.

Presumably, most of the anisotropy arises from texture (nonrandom crystallite orientations).

Anisotropy of the shear-mode internal friction Qq varied from 1.14 to 2.74, a surprisingly narrow range given the possibilities arising from changes in microstructure. During thermal

treatment, the C series steels showed decreasing Q"' anisotropy, probably reflecting dislocation

rearrangement.

For all five steels, all the usual quasiisotropic elastic stiffnesses B, E, G, Cl, lie within a band 2% below the values of pure Fe. Thus, to use these stiffnesses as a quality parameter

requires a measurement sensitivity of better than 0.1%, which we have in both the ARS and MOR

measurement methods.

As a function of aging treatment, the C series steels show a tight monotonic increase in

the Cjj, increasing about 1% from the underaged to overaged conditions. (See Figs. 2.3.1 and 2.)

The internal friction provides even larger promise for hardness correlations. For the C series

steels, where hardness increases about 3 digits on the Rockwell-A scale upon going from

underaged to peak aged, Qq decreases from about 13 to 8 and Q"g from 21 to 4. Presumably,

the decrease results from precipitation of small coherent Cu-rich particles, which pin dislocations,

which we assume make the largest contribution to Q'V (See Figs. 2.3.3 through 2.3.5.) We

measured both the strain and frequency dependences of Q'', both useful probes of the lattice

defects. Both results present surprises. (See Figs. 2.3.6 through 2.3.9.) The Q'' straindependence curve rises rapidly and reaches a saturation value. Figure 2.3.7 gives a possible

explanation. The Q'^ frequency dependence shows a negative slope (Fig. 2.3.8). At frequencies

below GOg, the Koehler-Granato-Lucke [2.3.1] model predicts a positive slope. This opens the

possibility of magnetic contributions to Q''. However, magnetomechanical hysteresis shows no

frequency dependence. Microeddy currents show Q'' proportional to frequency f. Macroeddy

currents show the same proportionality at low f and a f

''^^ dependence at high f. Thus, we fail to understand the negative slope dQ'Vdf.

21
Using MOR equipment at 150 kHz, we measured the Young modulus E and the

associated internal friction Q'' between 295 and 200 K, for the CA steel. (See Fig. 2.3.10.)

During cooling, the Young modulus showed a smooth monotonic increase with a slope higher

than for pure iron. We expect this from alloying effects and the fact that the Fe-Cu steel has

lower elastic stiffness than iron. The E-T curve showed no detectable irregularities upon cooling

through the ductile-brittle transition. The Q'' versus T curve showed a surprising result: a

relatively sharp peak near 265 K, significantly below the ductile-brittle onset. (See Fig. 2.3.1 1.)

If further measurements on other steels confirm this peak, then we speculate that it arises from

dislocations. One proposed explanation of the ductile-brittle transition is that the a/2<l 1 1> screw

dislocations become immobilized by defects such as interstitials, substitutionals, precipitates.

By collaborating with a non-NIST research group [2.3.2], we used a forced-vibration

torsional pendulum to measure Q'' at frequencies from 0.001 to 10 Hz (4 decades) for CA, C5,

C7 steels. (See Fig. 2.3.12.) Absolute Q"' values agreed well with ours measured at kHz to MHz

frequencies, and the results show the above-described V-shape behavior: CA>C7>C5. However,

the spectrum failed to show an expected peak near 0.1 Hz. After annealing the C5 specimen at 525 °C for 5 minutes, we found a strong internal-friction peak near 0.1 Hz. Thus, if, after thermal

treatment, these alloys sit at ambient temperature for a long time, then microstructural changes

eliminate the near -0.1 Hz peak in the mechanical-spectroscopy spectrum. This peak arises from

carbon-lattice interactions. But it could prove useful in studying copper-atom distribution

because the Cu-C interaction must differ from the Fe-C interaction.

We considered correlations with mechanical hardness. The C series steels stiffen

elastically with increasing hardness, but soften beyond the optimum-aging treatment. In the Csteel case, in going from hardness 56 to hardness 58, the shear modulus increases about 0.4%,

almost 10 times our measurement uncertainty of 0.05%. Thus, we can correlate elastic stiffness

and hardness, which relates in turn to yield strength, and to the ductile-brittle-transition

temperature. See Figs. 2.3.13 and 14.

Finally, we considered texture, which can alter drastically the elastic-coefficient values.

We assumed orthotropic macroscopic symmetry: x' = rolling direction, x^ = transverse direction,

x^ = plate-normal direction. From the measured Cy, we calculated the orientation-distribution

coefficients W.j^. We used these to prepare (1 10), (1 1 1), (200) pole figures for the CA, C5, C7

steels. (See Fig. 2.3.1.5.) As expected, the CA and C5 pole figures showed identical pole

distributions. The C7 pole figure differed slightly, suggesting the beginning of recrystallization,

which would alter both the macroscopic elastic stiffnesses Cy and their associated internal frictions

Qj]. All the pole figures show a hint of the usual b.c.c. rolling texture: (100) planes perpendicular

toxl

22
2.3.4 Conclusions

From these measurements on elastic stiffnesses C,j and internal frictions Q-], we draw the

following conclusions:

1

.

With thermal-treatment-induced microstructural changes, the C,j change is small,

(about one-half percent) but large compared to our measurement uncertainties.

2. Qj] shows a much larger change. For example, some of the Qi] are about 40% lower

for the C5 steel than the CA steel.

3. Although the relationship is not monotonic, both C,j and Q;] change with hardness, as expected from the presumed microstructural changes.

2.3.5 Recommendations

Based on the present results, we advocate the following further studies:

1

.

Assuming that Cu atom distribution affects ductile-to-brittle behavior, study a series of

Fe-Cu binary alloys free from complications of other alloying elements.

2. Further studies on internal friction, including especially the effects of temperature,

strain, and frequency.

3

.

Consider magnetic-field effects on reducing total internal friction.

4. Perform additional low-frequency torsional-pendulum measurements because this

method appears to provide the most-accurate internal-friction values.

5. Explore the possibility of getting useful information from the Snoek peak. Carbon

atoms provide a potent anisotropic strain field, which should change when other impurity atoms

such as copper occupy the lattice.

6. Conduct basic studies on the substitutional-interstitial internal-friction problem (mainly

Snoek peak). This would require a series of Fe-Cu-C alloys.

23
2.3.6 References

2.3.1 Lucke, K.; Granato, A.; Application of Dislocation Theory to Internal Friction Phenomena

at High Frequencies, J. Appl. Phys. 27:789; 1956.

2.3.2 Commercial instrument made available by Professor Wuttig of The University of

Maryland.

24
Table 2.3.1 Elastic coefficients and internal frictions of C steels.

CA

(underaged)

C5

(optimum-aged)

C7

(overaged)

Iron

p (g/cm^) 7.8414 7.8558 7.8534 7.872

Cii (GPa) = 272.96±0.46 273.44±0.57 275.38±0.72

C22 (GPa) = 275.37±0.30 275.75+0.39 276.16+0.44

C33 (GPa) = 272.13±0.73 273.48+0.90 273.26+0.98

C12 (GPa) = 105.77±0.36 105.44+0.44 106.52+0.81

Cl3 (GPa) = 1 1 1 .47±0.64 111.90±0.78 112.06±0.82

C23 (GPa) = 107.93±0.60 108.03+0.75 108.89±0.51

C44 (GPa) = 78.725±0.000 78.944+0.000 79.804 ±0.008

C55 (GPa) = 82.371+0.008 82.684+0.008 82.975 ±0.008 Cee (GPa) = 77.571±0.023 77.955±0.023 78.414 ±0.024

Sn (TPa-^) _ 4.739 4.727 4.695

S22 (TPa-^) = 4.632 4.617 4.628

S33 (TPa-^) = 4.791 4.767 4.776

S.2 (TPa-^) = -1.255 -1.242 -1.248

Si3 (TPa-^) = -1.444 -1 .444 -1.428

S23 (TPa-^) = -1.323 -1.316 -1.332

S44 (TPa-^) = 12.70 12.67 12.53

S55 (TPa-^) = 12.14 12.09 12.05

See (TPa-^) = 12.89 12.83 12.75

276.85

111.97

82.44

4.709

-1.356

12.13

B (GPa) 163.42 163.71 164.41 166.93

E, (GPa) = 211.00 211.56 212.99

E2 (GPa) = 215.89 216.60 216.09

E3 (GPa) = 208.73 209.77 209.38

Q-^2(10-')* = 9.94 6.40 7.26

V12 = 0.2647 0.2627 0.2658

Vis = 0.3047 0.3054 0.3042

V23 = 0.2857 0.2850 0.2879

V21 = 0.2708 0.2690 0.2696

V31 = 0.3014 0.3028 0.2990

V32 = 0.2762 0.2760 0.2790

212.36

0.2880

Xi = rolling direction; X2 = transverse direction; X3 = plate normal

*Measured Young-modulus-mode internal friction in direction 2.

25
Table 2.3.1 (continued)

CA C5 C7

(underaged) (optimum-aged) (overaged)

Q\(W^)* = 8.47 9.53 7.99

0-^2 (10 M* = 5.12 3.38 3.75

Q\ dO'')* = 15.4 4.41 3.64

Q\i (10^) — 8.64 9.68 10.8

Q'22(10M = 15.4 2.70 5.58

Q'aadO^) = 19.8 2.57 4.95

QV(IO') = 19.5 7.06 11.9

Q'ssdO") = 10.3 10.9 11.0

Q'eedO'') = 9.58 10.0 10.4

Q^B (10 ») ^ 20.88 4.32 8.88

Q^6 (10^) = 13.30 7.99 8.45

Q-^E (10=) = 14.37 7.47 8.51

Q-\^.(10*) = 9.94 6.40 7.26

X, = rolling direction; Xj = transverse direction; Xg = plate normal

'Averaged Young-modulus-mode internal friction.

i

26
Table 2.3.2 Elastic coefficients and internal frictions of 1-2 and W1 steels.

1-2 w1 Iron

p (g/cm^) = 7.8532 7.8339 7.872

Cii (GPa) 274.42 273.54

C22 (GPa) 274.44 274.50

C33 (GPa) 274.49 273.40

Ci2(GPa) 110.34 110.19

Ci3(GPa) 110.54 109.30

C23 (GPa) 110.54 109.56

C44 (GPa) 81.987 82.007

C55 (GPa) 82.011 82.234

C66 (GPa) 82.001 82.202

Sn (TPa-^) 4.738 4.745

S22 (TPa-^) 4.738 4.730

S33 (TPa-^) = 4.740 4.738

Si2(TPa-^) -1.356 -1.366

Si3(TPa-M -1.362 -1.350

S23 (TPa-^) -1.362 -1.349

S44 (TPa-^) 12.20 12.19

S55 (TPa-^) = 12.19 12.16

See (TPa-^) = 12.20 12.17

276.85

111.97

82.44

4.709

-1.356

12.13

B (GPa) = 165.13 164.39 166.93

E-i (GPa) = 211.07 210.74 212.36

E2 (GPa) = 211.08 211.42

E3 (GPa) = 210.97 211.07

V12 = 0.2863 0.2879 0.2880

V13 = 0.2874 0.2844

V23 = 0.2874 0.2853

V21 = 0.2863 0.2889

V31 = 0.2873 0.2848

V32 = 0.2873 0.2848

x-] = rolling direction; X2 = transverse direction; X3 = plate normal

27
Table 2.3.2 (continued)

1-2 W1

Q\, (10^) = 6.564 5.568

Q'zzdO^) = 5.599 6.356

Q^adO-S) = 11.46 0.499

Q^44(10') ^ 10.23 13.28

Q^sdO-^) = 6.813 6.363

Q-^66(10-') = 5.396 12.70

q^b(io-') _ 9.369 3.409

Q^G (10-5) = 6.735 8.557

Q-\ (10 5) = 7.109 7.824

Xi = rolling direction; Xj =transverse direction; X3 = plate normal

Table 2.3.3 Elastic-stiffness anisotropy defined as ratio of maximum to

minimum shear-mode elastic-stiffness coefficients (C44, C55, Cge). Also, we

show a similarly defined anisotropy for the internal friction Q'^.j.

Steel C|j anisotropy* Q \j anisotropy

CA 1.062 2.04

C5 1.061 1.54

C7 1.058 1.14

1-2 1.000 1.89

W1 1.003 2.09

*For an isotropic nnaterial, this value equals 1.000. For polycrystal iron, the maximunnpossible C|j anisotropy equals 2C44/(Cn-Ci2) = 2.407.

28

1
1.005

O

LL

O

o

1 .000 -

0.995 -

0.990 -

0.985 -

0.980 -

0.975

Various Steels

Fig. 2.3.1 Polycrystal (quasiisotropic) elastic stiffnesses for the five steels compared with iron.

B = bulk modulus, Cl = longitudinal modulus, E = Young modulus, G = shear modulus, v =

Poisson ratio. Precipitation of atoms from solid solution increases stiffness.

29
1.010

1 .005 _

1.000_

o

0.995 _

0.990_

0.985 _

0.980_

0.975

CA C5 C7 W1 1-2 Fe

Various Steels

Fig. 2.3.2 Companion to Fig. 2.3.1. Derived monocrystal elastic stiffnesses. Ci,, C12, C44 = usual

Voigt elastic coefficients, and C = !/2(Ci, - C12). Anisotropic stiffness behaves similar to quasiisotropic case.

30
0.6 -

X 0.4-

0.0

CA C5 C7 W1 1-2 Fe

Various Steels

Fig. 2.3.3 Companion to Fig. 2.3.1. Shows internal friction Q' for various quasiisotropic modes.

In the C steels, probably loss of precipitate coherency causes V-shape.

31
li.

a

Various Steels

Fig. 2.3.4 Companion to Fig. 2.3.1 and 2.3.3. Shows Q"' associated with the three anisotropic

shear moduli. Different behavior for the 44, 55, 66 cases suggests anisotropic dislocation

geometry.

32
<

o

I g

0.2

CA C5 C7

Fig. 2.3.5 Relative Q"' values for C steels. Ej refers to MOR measurements, E to ARS

measurement. Two independent measurements confirm V-shape behavior.

33
6

15

14

13

12-

11 -

10-

9-

5-

' ' ' I D

I I I I

D

O

-O

J I I I I I I

' '

'44

O

o o

C33 + 3.6Cgg

AA2 SteelI I J I L. I I I 0.0 0.2 0.4 0.6 0.8 1.0 1.2

Strain (lO"^)

Fig. 2.3.6 Q' strain dependence. Both C44 and €^6 represent shear modes. The AA2 steel is a

lower-copper-content companion to the CA steel. Saturation above about 0.3 x lO"* departs from

usual Koehler-Granato-Liicke model.

34
Hypotheses to explain previous figure.

crit

Region 1 vibrating dislocation

viscoelastic

reversible

Region 2

moving dislocation

microplastic

irreversible

Fig. 2.3.7 Companion to fig. 2.3.6. Possible interpretation of the Q"' strain relationship. Low £,

means high ductility.

35
0.4

0.3

CO

I o

1 1 1 I

0.2

0.0

1 I 1 1 1 1 r- CA Steel

100 200 300

Frequency (kHz)

400

Fig. 2.3.8 Q"' - frequency dependence. Surprising negative slope disagrees with usual dislocation

and magnetic mechanisms. Scatter arises from various deformation modes.

36
1E-7

1E-8

C

"(5

(0

a

lE-9 -

1E-10

T—I—I—J—[—I—r—I—I—

I—I—J—I—I—

I—I—i-T—I—

r

I I I I AA2 Steel

8

a

o

C33 + 3.6Cg6

C44

J I I I L_l L-J I I I L-J I L J I I I I I I L

oo

J U-J L. 10 20 30 40

[Strain (10-^)] -4\t1

£ = 2*10-

50

8 = 2*10-^

60

Fig. 2.3.9 Granato-Liicke diagram. The Koehler-Granato-Liicke model predicts linear behavior.

Departure from linearity at low strains remains to be explained; it occurs in many materials. The

AA2 steel is a lower-copper-content companion to the CA steel.

37
1.04

1.03

lO

LU

LU

1.02

1.01

i.ooi 1 150

I I I

I

I I I I CA Steel

TEMPERATURE (K)

Fig. 2.3.10 Temperature variation of CA-steel Young modulus compared with iron. Measured

using MOR method. Upon cooling through the ductile-brittle transition, no significant irregularity

appears. (Data for iron, see Ref. [2.3.3].)

38
in

CM

O

O

5

4

1 1 1 1 1

. 1 1

-t \ '

'. CA Steel

f\ CO _ 3

o \

9 • . P

2

- - o-

•• cP

•. - o

6 \

^o

""o^^ ^ "o^' 1 ° o

o ^^o

o P

o

n

150 200 250

TEMPERATURE (K)

300

Fig. 2.3.1 1 Companion to Fig. 2.3.10. Peak in internal friction implies a relaxation process near 270 K, close to the ductile-brittle transition temperature.

39
<o

c

o

c

0)c

Q.Q008

0.001 0.01 0.1 Frequency (Hz)10

Fig. 2.3.12 Snoek peak of C5 steel. Measured at a non-NIST laboratory using a forced-vibration

torsion pendulum. The Snoek peak arises from carbon-atom reorientation under stress. Replacing Fe-lattice atoms with other atoms (a ternary solute addition such as Cu) should change

this peak: broadening, frequency shift, peak height, subpeaks.

40
0.984

0.982

0)

UL

O

O

0.980 -

0.978

1 1 1 1 I

1

1 I

1 r C Steels

54 56

DG

CA C5

J I I I I I I I I 58

Hardness (R.)

Fig. 2.3.13 Elastic-stiffness change with hardness.

60

41
LT)

I o

zo 1 1 1 1

I 1

1 f-- .

.

> C Steels •

-

- O

• 20 \ - \

\

\ - \ - 15 \ -

D \

10 -

\ "

- \ DG

5

- \ :

- Ob

CA C5

1 1 1 1 1 54 56 58

Hardness (R.)

60

Fig. 2.3.14 Companion to Fig. 2.3.13. Q"' change with hardness.

42
CA specimen

C5 specimen

C7 specimen

Fig. 2.3.15 Principal pole figures (100, 110, 1 11) derived from acoustic measurements, which

yield the orientation-distribution coefficients W,j;^. These reveal texture (nonrandom-crystalliteorientation distribution).

43
2.4 X-RAY DIFFRACTION —Davor Balzar

2.4.1 Methodology

Because the X-ray strain gauge is so small, the diffraction method is sensitive to both

macrostress (generally caused by any applied or residual macroscopic elastic and plastic

deformation of material) and microstress (for instance, caused by elastic incompatibility of

different grains in a single-phase material, different plastic deformation or thermal expansion

coefficients of phases in multi-phase materials). Terminology used in the literature divides

residual stresses into three kinds: Kind I - acting over macroscopic regions; Kind II - acting over

sub-millimeter dimensions (such as grains); and Kind III - acting over sub-grain dimensions at the

crystalline lattice level. The first two categories shift diffraction lines because the deformation is homogeneous throughout the sampling volume. Residual stresses of Kind III are visible mostly in

line broadening, caused by localized inhomogeneous deformation associated with structural

defects. Analysis of diffraction line-broadening can quantify both size-related defects (such as stacking and twin faults, low-angle and high-angle boundaries) and strain-related defects (namely

dislocations, point defects, and distortions of the host lattice in the vicinity of coherent

precipitates). Diffraction methods measure directly strain and not stress. Conversion between the

two is provided by the way of macroscopic elastic constants. However, because the Kind III stress is localized on so small a scale (in the 1 to 100 nm range), the simple linear elasticcontinuum approximation may not be valid. Therefore, we have chosen here to report only on

the measured quantity: the strain.

2.4.2 Experimental Technique

Specimens CA, C5, and C7 were prepared by polishing and etching one surface. The

average x-ray sampling volume was about 15 mm x 3 mm x 5 |um. Measurements were made on a

horizontal goniometer with CuKa, 2 radiation in a high-resolution setup. The complete diffraction

patterns were collected (six Bragg reflections for the ferrite structure) and showed peaks from

only the a-Fe phase plus a few very weak peaks of unknown origin. It is not possible to distinguish between the peaks from the main ferrite phase present and bainite, but because no

carbide peaks were observed, the contribution of bainite should be minimal. The diffraction

patterns exhibited relatively strong texture, but this varied greatly among the specimens and may

depend on how the specimens were cut. (See Fig. 2.3.15 for pole figure information.)

2.4.3 Results

The diffraction patterns were modeled with a Rietveld-refinement program. The refined

lattice parameters did not differ significantly, which indicates no change of hydrostatic strain or chemical composition between the specimens. Line broadening was moderate in all three

specimens, but it did differ significantly, especially between C5 and C7 specimens. Preliminary

diffraction-line broadening analysis indicated an insignificant effect of anisotropy and only the

isotropic line-broadening parameters were refined by the Rietveld program. Coherent-domain size

44
and strain were estimated from refined diffraction-line full-widths-at-half maximum, which were

corrected for instrumental line broadening. The refined domain sizes were larger than about 300

nm (which indicates a very small contribution to line broadening), and increased monotonically

but not significantly from CA to C5 to C7. Strain-induced broadening was dominant, although

the magnitude was not large (less than 0.1%). Strain values are presented in Table 2.4.1 for the

C-series, 1-2, and Wl specimens, and are plotted as a function of hardness in Fig. 2.4.1. A good

linear correlation is evident. Different offsets for the A-series and C-series of specimens may

indicate somewhat different physical mechanisms of broadening in the two alloy types. Fig. 2.4.2

shows a good correlation between strain and ultimate tensile strength, as well.

2.4.4 Physical Model

Because it is difficult to prevent the formation of copper-rich precipitates, the CA

specimen probably contains a relatively large number density of such precipitates. However, they

are probably significantly smaller than those in the C5 specimen. This is confirmed by relatively

higher strain and hardness of the CA compared to the C7 specimen. There is a sharp drop in strain

between the C5 and C7 specimens. A possible explanation for this is that the coherent copperrich precipitates, which are of optimum size in the C5 specimen, coarsen due to higher and

prolonged annealing temperature (overaging), thus becoming incoherent with the matrix, which in

turn relieves the matrix strain. Moreover, the effect may be accompanied by a simultaneous

decrease in dislocation density because coarser precipitates do not pin dislocations so efficiently

and the material has been subjected to an annealed higher temperature. The reason that a drop in

hardness and ultimate tensile strength is not accompanied by a decrease in the proportional limit

(see Fig. 2.2.2) may simply be because there is an overall lower dislocation density in the C7

material compared to the C5 specimen.

Table 2.4.1 Strain deduced from diffraction line-broadening analysis.

Specimen Strain ( 1 0'"*) CA 6.9(5)

C5 7.4(1)

C7 5.3(5)

1-2 7.3(1)

Wl 7.7(1)

45
0.0015 -

^ 0.0010

0.0005 -

AA

-

"""

A6 A5

* *

CA

C5

AA-2 C7 I

*

i

< 1 1 50 55

Hardness (R^)

60

Fig. 2.4.1 Relationship between internal lattice strain and hardness for the A710 steel (CA, C5,

and C7) and the 1508 steel (AA, AA-2, A5 and A6).

U.UUIU

, 0.0009 -

0.0008 - CA

C5

0.0007 -

II .s . ^ 0.0006 - C7

0.0005 -

II

0.0004

ft ftftfti

-

1 .„ 1 J

1 1 550 600 650 700

Ultimate Tensile Strength (MPa)

750

Fig. 2.4.2 Relationship between internal lattice strain and the ultimate tensile strength for the A

710 alloy in the overaged (C7), the solution treated (CA), and the peak-aged (C5) condition.

46
2.5 NONLINEAR ULTRASONIC PROPERTIES—Donna Hurley

2.5.1 Background

When an ultrasonic wave at a frequency o) and amplitude A, propagates through a solid,

nonlinear terms in the force-distance relationship that describes the interatomic forces cause the

generation of a wave at frequency of 2o) (a second harmonic). The amplitude of this second

harmonic wave Aj increases with distance and with the square of the amplitude A^ . By

measuring A2 relative to A,, it is possible to deduce the value of a parameter P that is related to the nonlinear terms in the interatomic force relationship and is, therefore, a physical property of

the solid. Harmonic generation experiments were performed on eight steel samples in order to determine their nonlinear ultrasonic parameters and to establish their relationship to the

microstructures. Three of the samples contained 1.13% copper (the "C" series) and three

contained 0.35 mass percent copper (the "A" series). The remaining two samples were extracted

from the SNUPPS reactor vessel and consisted of A533B steel from the region adjacent to a weld (sample 1-2) and from the weld itself (sample Wl). The "C" and "A" samples had received

different heat treatments in order to vary their mechanical hardness, as described in Section 2.1 of

this report.

2.5.2 Experiments

In the experiments, a tone burst of finite amplitude and at a frequency of 9.8 MHz was

launched with a piezoelectric transducer bonded to one side of the specimen. The absolute

displacement amplitudes of the wave components at o) (fundamental) and 2(o (second harmonic) were measured on the other side of the specimen using an infrared, path-stabilized Michelson

interferometer. Experimental techniques for determining (3 in this way have been described in

detail elsewhere [2.5.1]. Values for P were calculated from the ratio of the fundamental and

second-harmonic wave amplitudes together with values for (o, the longitudinal velocity v^ and the

sample thickness. (The velocities were measured by K.W. Hollman using a high-precision pulseecho superposition technique [2.5.2].) Physically, p is defined in terms of the second and thirdorder elastic stiffness moduli. For the experiments performed here, p = -(3 + C, , /C, ,).

2.5.3 Results

The experimental results are summarized in Table 2.5.1. It lists both the absolute value of

P and its value P' relative to that for the C7 specimen. By using relative values, it is possible to eliminate common sources of measurement uncertainty and thus reduce the relative uncertainty to ± 2%. Also given are values for the residual strain as determined by x-ray diffraction linebroadening experiments (discussed in Section 2.4 of this report).

47
Table 2.5.1 Compilation of measurements associated with nonlinear elasticity.

Sample Hardness (HRA) P

p' (relative) VL(m/s) 8(10-^)

C7 53.2±0.5 4.81±0.29 1 5906±10 5.3±0.5

CA 55.6±0.5 5.16±0.31 1.07±0.02 5885±10 6.9±0.5

C5 58.5±0.5 5.36±0.32 1.11±0.02 5898±10 7.4±0.1

AA2 50.0±0.5 5.64±0.34 1.17±0.02 5903±10 N/A

A6 53.1±0.5 5.42±0.33 1.13±0.02 5906±10 9.1±0.1

A5 54.3±0.8 5.66±0.34 1.18±0.02 5906±10 9.3±0.2

1-2 53.3±0.8 5.30±0.32 1.10±0.02 5912±30 7.3±0.1

Wl 57.0±0.5 6.56±0.39 1.37±0.02 5917±30 7.7±0.1

2.5.4 Discussion

The results for the three C series specimens are plotted in Fig. 2.5.1 as a function of

hardness. This figure reveals that P changed markedly with sample hardness and increased by

11% over a hardness change of about 10%. We think that this behavior is due to changes in the

copper precipitate strain configuration with tempering—precipitate formation and growth of

coherency strain, then gradual loss of coherency through coarsening. At peak hardness (C5

sample), the precipitate is fully coherent and the surrounding region contains a great deal of strain. In overaged material (C7 sample), the precipitate becomes incoherent, relieving the strain. The

nonlinear ultrasonic experiments measure an average strain in the bulk of the sample and,

therefore, they should be sensitive to the relative state of precipitate coherency [2.5.3]. A

comparison of the nonlinearity parameter P with the x-ray strains, given in Figure 2.5.2, shows an

excellent correlation. Since nonlinear ultrasonic experiments probe bulk, and not just near-surface,

properties (as in the case of the x-ray experiments), these results offer promise for nondestructive

characterization of precipitate strain. Further experiments are being planned to obtain a more

systematic understanding of this behavior.

The correlation between strain and the nonlinear parameter P for the C series leads us to ask whether the remaining data exhibit similar trends. Accordingly, all of the experimental results

for p as a function of strain are plotted in Fig. 2.5.3. This figure reveals that P increases

approximately linearly with strain for all of the samples measured, including the A533B reactor

steel, except for the weld specimen (Wl). In the weld specimen, precipitate strain is probably not

the only microstructural feature that contributes to harmonic generation. Therefore, the weld

48
specimen results serve as a reminder that the absolute value of P cannot be interpreted as a direct

reading of the microstrain, although the relative values may be useful to assess the relative state of

strain.

2.5.5 References

2.5.1 Hurley, D.C.; Fortunko, CM.; Determination of the nonlinear ultrasonic parameter P

using a Michelson interferometer, Meas. Sci. Tech. 8: 634; 1997.

2.5.2 Renken, M.R.; Kim, S.; Hollman, K.W.; Fortunko, CM.; Accuracy of Resonant

Ultrasound Spectroscopy: an Experimental Study, Proc. 1997 IEEE Ultrason. Symp., Vol. 1, 97CH36118, ISSN 1051-0117, p. 483.

2.5.3 Cantrell, J.H.;Yost, W.T.; Effect of precipitate coherency strains on acoustic harmonic

generation, J. Appl. Phys. 81:2957; 1997.

49
1.20

1.10-

1.00 -

0.90

54 56

hardness (HRA)

Fig. 2.5.1 Relationship between the nonlinearity parameter P and the Rockwell A Hardness of the

A710 alloy after heat treatment to different hardness levels.

ea

^.^{J 1 1 1

1.10 -

y

y ^yy

y _

r y I. y

y

y

y ^

1.00 ' :^ y

y

y

y

y

y

y

y

y

y

y

K •

0.90 1 1 1 4 5 6 7 8

e(IO^)

Fig. 2.5.2 Relationship between the nonlinearity parameter P and the internal lattice strain

deduced from X-ray diffraction line broadening.

50
6-

5-

1^

6 8

e(10-^)

10

Fig. 2.5.3 Relationship between the nonUnearity parameter and the internal lattice strain for the

A710 alloy (squares), the 1508 alloy (circles) and the A533B alloys (triangles).

51
2.6 MAGNETIC PROPERTIES —Fred Fickett

2.6.1 Introduction

Magnetic measurements for detecting embrittlement of the reactor steels are not new. The

interaction of moving domain walls with lattice defects of various sorts is well known as a pinning

mechanism that shows up as dramatic changes in the magnetic coercivity with relatively small

changes in defect type or concentration. Other magnetic parameters, such as permeability,

remanence, and saturation magnetization, also offer the potential for providing information on

specific defect or compositional effects, although in all cases the interactions are very complex.

Magnetic measurements offer one of the best chances for a detection system that can be used in

the actual pressure vessel environment where the detection must be done through a relatively

large thickness of stainless steel cladding, which normally is a normiagnetic material. In this part of

the project, we are looking at steels of various compositions with conventional magnetic

measurement systems. These data will form the basis for interpretation of the necessarily less

sophisticated measurements to be made in the field.

2.6.2 Literature Review

The existing literature on magnetic properties of reactor steels was reviewed. Much of the

literature is embodied in a series of progress reports to the NRC over a large number of years as well as conference papers based on those reports: a lot of paper and not many data. The final

report was published just last year and is available only on microfiche. Measurement techniques

used were mostly of the large-sample B-H looper type. At best, existing data on the effect of the

precipitates on magnetic properties are inconclusive. Interestingly, it seems that no precise

magnetic measurements, such as can be made with the vibrating-sample magnetometer, have yet

been done.

2.6.3 Vibrating Sample Magnetometer (VSM) Measurements

The VSM is a laboratory instrument that will almost certainly not be adaptable to field use. However, it gives the best direct data on the magnetic properties of the bulk steel. Specifically, it is capable of measuring the entire hysteresis loop of the sample, from which all of

the static magnetic parameters can be obtained. These data will be needed as baseline or input to actual field measurements as well as for the detailed understanding of the magnetic behavior of

the steels. At this stage of the investigation, the instrument allows us to answer many questions

quickly and easily.

2.6.4 Apparatus

The VSM concept is shown in the Fig. 2.6.1 . The steel sample is vibrated in an applied

(variable) magnetic field and the moment induced by the field is detected by a set of pickup coils

next to the sample. Clever design of the pickup coil and sophisticated modern electronics allow

52
various sources of error to be corrected out so that the final readout is a voltage proportional to the magnetic moment of the sample. The VSM borrowed for the tests reported here had been

optimized for very low level magnetic measurements and our samples are at quite the other end of

the scale. Test runs on nickel and steel samples indicated that the apparatus and the associated

data acquisition programs needed modification to accurately measure materials of high magnetic

moment. The VSM coil set and the location of the applied field magnetic Hall probe pickup were

modified to avoid interactions between the sample and the probe. A VSM always has to be

calibrated with a sample of known moment, preferably of the same size and shape as the samples

to be measured. In our experiments, this was done with a pure nickel sample.

2.6.5 Samples and Sample Preparation

In all cases, the steel samples measured were in the form of rods 2 mm in diameter by 20

mm in length cut by electric-discharge machining (EDM) from the bulk material with their length

in the plane of the sheet, essentially at the center of the thickness. Specimens were prepared from

three steels, the A710 steels (CA, C5, C7), the 1508 steels (AA, A5, A6), and A533B (1-2) and a welded A533B(W1).

The calibration sample was cut from a 2 mm rod of pure nickel (99.9945% Ni). Because

of the purity, the saturation moment was easily calculated and it is that number that calibrates the

VSM at the beginning of each set of measurements.

2.6.6 Data and Calibration Runs

As mentioned above, the calibration run with the nickel sample was made frequently,

necessitated in large part because the instrument was used by a number of people in the course of

a week and the calibration quickly degrades if major changes have been made to the apparatus.

Although, the calibration data consist of only the saturation magnetization of the sample,

complete hysteresis loops, as described below, were always run.

For each of the steel samples, two types of measurements were made: The first was a

hysteresis loop extending to 240 kA/m (3000 Oe) applied field and starting with magnetization to -240 KA/m (-3000 Oe). Each of the 200 points on these loops is a dc measurement, i.e. the

applied field was stationary during the actual moment measurement. This curve was used to determine the coercivity. The second curve was started at zero field with a demagnetized sample

and taken to 368 kA/m (4600 Oe) on a coarser grid of 75 points. This curve was used to determine the saturation magnetization. The measurements described above were completed on

one sample of each of the steels for a total of 16 data runs.

2.6.7 Data Analysis

To date, analysis of the data to determine the saturation magnetization and the coercivity

of the steel samples has been completed. In essence this involves plotting the portion of the loop

53
Table 2.6.1 Values for the magnetic parameters measured with the VSM.

Sampe Condition HsCOe) H,(A/m) ^IqMs

CA Solution treated 8.39 668 2.03

C5 Peak aged 8.07 642 2.07

C7 Overaged 6.10 485 2.06

W-1 Welded A533B 8.95 712 2.05

1-2 Bulk A533B 6.84 544 2.06

in the appropriate region and doing a linear fit to the data. Figure 2.6.2 shows this process for

determining the coercivity of the C5 steel. The magnetizaton at saturation is easily extracted from

the virgin field curve, although the value at 368 kA/m (4600 Oe) is little changed fi^om that seen

here at 240 kA/m (3000 Oe). A summary of the data is given in Table 2.6. 1 . Note that the A

steels are not included; fiirther assessment of their appropriateness for this project led us to drop

them from the sample set. Remanence data were also analyzed, but for these steels they add

nothing that is not already seen in the coercivity. The curves were not analyzed for permeability

values at this time. That analysis requires a fairly lengthy process of correcting the curves for the

demagnetizing field of the sample. If later investigation indicates that it might be a valuable

parameter, it will be extracted then.

2.6.8 Low Temperature VSM Experiments

The coercivity He and the saturation magnetization, ^qM^, were measured at room

temperature and on cooling to liquid nitrogen temperature (76 K) on new samples of A533B

taken from the base metal (sample 1-2) and from the weld (sample Wl). The results are shown in

Table 2.6.2. In all cases, the effect is essentially reversible. Based on these limited data, we

concluded that we are not seeing a transition of any sort, but that the observed increase is an

effect of reduced thermal energy which, at the higher temperature, hinders the alignment of

moments, lowering M^, and helps the motion of domain walls, lowering H^,. Complete proof

would require an experiment using a larger number of temperature points. H^, data on longer

samples from earlier measurements which were cut from different regions and prepared slightly

differently from the above gave values of H, = 6.84 (1-2) and H^= 8.95 (Wl).

54
Table 2.6.2. Values of the coercivity H^ and saturation magnetization M, for the A533B

steels and the nickel standard at room temperature (two readings) and at the

temperature of liquid nitrogen. (The samples used were 2 mm dia. x 10 mm

long.)

Parameter T 1-2 Wl Nistd

Room temp. 571 (7.18) 759 (9.54) 1990(25.0)

H, in A/m (Oe)

(2-point method) Liquid Nj 602 (7.56) 844(10.60) 2316(29.1)

Room temp. 571 (7.18) 761 (9.56) 1997(25.1)

Room temp. 59.2 59.2 18.5

M, (mV)

(not calibrated) Liquid Nj 61.5 61.7 19.5

Room temp. 59.2 57.9 18.4

2.6.9 Results and Conclusions

Preliminary analysis of the data for the C- series steel indicates a dependence of coercivity on sample treatment, but no dependence of the saturation magnetization. This latter observation is to be expected given that the composition of the steel is not changed dramatically by the heat

treatments. The coercivity variation does not track with the hardness, which is taken as the measure of embrittlement for these experiments. Again, this is not surprising given that a number

of changes in the microstructure, defect content, and arrangement are occuring in this material.

There is still reason to believe that, when applied to the actual problem where the confounding

variables are minimized, the magnetic measurements may well provide excellent diagnostics.

There are several clear directions for the future of these measurements: First, we need to evaluate a number of samples from a given lot of steel with more tight control on their

preparation, especially the EDM operation and subsequent cleanup. Second, we need to investigate the possibility of measuring a sample of the irradiated steel in our laboratory. The

samples are sufficiently small that this may be possible. If even smaller samples were necessary, our SQUID magnetometer systems could measure milligram quantities, albeit with a bit more

difficulty than that encountered with the larger samples in the VSM. The major problem will be in

getting the appropriate samples prepared—almost certainly a hot-cell EDM system would be

needed.

55
2.6.10 Other Methods

B-H Looper. A number of techniques that we think may be used on samples of the

actual steel (unirradiated, but with cladding in place) make use of a large transformer core device

that could be called a B-H looper. We have built a first prototype of the device and made

preliminary dc tests (ac and dc are possible) with multiple field detector locations and have

demonstrated its ability to see a signal from the steel through the cladding. This is a long way

from a usable instrument, but we are encouraged by the results to date.

Bitter Techniques. Early in the project, we attempted some very preliminary

measurements on two early A710 samples (not the C samples) using the Bitter technique to look

for domain structure. We had no luck bringing the domains out either with a very fine oxide

powder or with a more conventional iron powder suspension. We probably need a smoother

surface prepared by electropolishing or chemical polishing and measured relatively soon after

preparation to avoid surface oxidation. This was only a secondary experiment and, thus, has not

been pursued further.

AFM/MFM Studies. Early in our thinking about the problem, we considered how to use

the atomic force microscopy (AFM), and maybe magnetic force microscopy (MFM), capability of

our scanning probe microscope to image precipitates with sizes in the 10 nm range. There is reason to believe that it can be done ifwe canfind the things and if the sample surface can be

prepared such that they remain there and don't get ripped out or flattened by the surface

preparation. No experiments were tried, but it remains a possible technique to be evaluated in

more detail.

Ion Implantation Studies. We have been involved in recent years in a project studying

nitrogen-ion-implanted stainless steel surfaces implanted with N ions. We can also implant Cu into

steel, and it occurred to us that this might make a reasonable model material for the irradiated

steel. Initially, two samples of Fe were implanted with Cu ions. These were very heavily implanted

(40% Cu), but in a very thin layer (-20 nm). It is possible that this may be a way to make a (very

small) sample with the right amount of copper. Certainly one can heat these samples to encourage

copper migration from the surface layer. New samples with better surface preparation and 1 to

1 0% implanted copper have been prepared and preliminary magnetic force microscope and atomic

force microscope images have been made. As usual, the data are confusing, but sufficiently

interesting to consider fiirther study under NIST funding.

56
Fig. 2.6. 1 Schematic diagram of a vibrating sample magnetometer (VSM).

C5 Steel Hysteresis

E

a>co

IBN

»

Ca

-100

•4000 -3000 -2000 -1000

Coercivity Determination

Applied Field, Oe

•20 20 Applied Field, Oe

Fig. 2.6.2 (Left) Example of the hysteresis curve observed with the VSM for alloy A710 in the

peak-aged condition. (Right) Magnified hysteresis curve near H = B = for determination of He.

57
2.7 MICROMAGNETIC TECHNIQUES — Brian Igarashi

2.7.1. Development of a Noncontacting Measurement of Magnetostriction

The objectives of the micromagnetics experiments are twofold: (1) demonstrating that

magnetostriction (the change of length of a ferromagnetic material that accompanies

magnetization) and the transverse incremental permeability are sensitive to the formation of Curich precipitates (CRPs) in carbon steel and (2) devising a noncontacting technique for measuring

magnetostriction. Such a noncontacting technique would be needed in the field, since the portion

of RPV carbon steel that is the target of inspection lies underneath stainless steel cladding.

Results of investigations into the first objective are discussed in individual sections below, but the

way by which the second objective is fulfilled may be described briefly as follows.

Magnetostriction of a steel sample is ascertained by measuring the amplitude of ultrasonic

waves generated by magnetostriction. The noncontacting aspect of the measurement is afforded

by use of electromagnetic acoustic transducers (EMATs) to generate and detect ultrasound.

Thompson [2.7.1-2.7.3] has shown that shear horizontal (SH) plate waves can be generated

magnetostrictively with the arrangement shown in Figure 2.7. 1 . A static magnetic field Hq

establishes an axis of tension due to magnetostriction. (Compression is regarded here as a

negative tension.) The ac current in the meander coil creates an oscillating field H^^ that adds

perpendicularly to the static field and, in effect, rotates the axis of tension in the surface plane,

thus, creating a shear strain.

According to Thompson, the amplitude of the SH wave obeys the relationship,

^oc5h xy

dHc,

(1)

where e^y is the magnetostrictive shear strain induced by Hq and H^^ (the derivative is referred to as magnetostrictive strain coefficient), h is the magnitude of H^^ at the sample surface, 6 is the skin

depth , Ht is the transverse incremental permeability, a is the conductivity, and o) is the angular

frequency. In near-saturation static fields, where the axis of magnetostrictive tension rotates nearly parallel with the net magnetic field,

^Exy 3 \

dH^ 2Hn* (2)

'CD

where X is the magnetostriction coefficient (the longitudinal strain due to Hq). At low Hq, this

expression gives an overestimate of the derivative, but we [2.7.4] recently have shown that

modelling of the wave amplitude is improved somewhat by using

58
Bexy 3 X UtH

an,., 'CO 2 H0)

HtHCO

CO B ;

CO MI

Br

CO

(3)

The range of validity of this expression extends to lower Hq than eq (2), because the axis of

magnetostrictive tension does not rotate parallel with the net magnetic field at low Hg.

Substitution of eq (3) into eq (1) yields better predictions of variations of wave amplitude at low

Hq, although relative values of the wave amplitude at low Hg are overestimated.

Estimates of | A, | are obtained by substituting measured wave amplitudes into the following

expression derived from eqs (1) and (3):

W«^Vi^

B, \itn

^2

CO l^^tHCO Br

Br \itnCO ^itHCO Br

(4)

This expression is valid primarily at high Hq, and that relative values of the estimated | A,| tend to be underestimated at low Hq . We experimented with a variety of methods for measuring the amplitude of SH waves

generated by EMATs (specifically, the fiindamental SHq mode). In initial experiments, we

attempted to infer the wave amplitude from the amplitude of standing waves set up between

parallel sides of a sample. However, it was discovered that the EMAT excited unwanted standing

waves that interfered with the one being monitored. Although this problem could be ameliorated

by sticking clay to the sides of the sample to dampen out unwanted standing waves,

measurements were judged to still contain an uncertain error. For this reason, this approach was

abandoned in favor of a pulse-echo technique that directly measures the amplitude of the traveling

wave generated by the EMAT.

One configuration of the pulse-echo measurement is depicted in Fig. 2.7.1 (Setup 1). A

meander-coil EMAT is used to generate SHq waves of 4 mm wavelength. A gated amplifier

issues pulses consisting of three cycles of sinusoids at 791 kHz, and the amplifier output is adjusted continuously to maintain a constant EMAT peak current of 5 A. The amplitude of a chosen echo is measured with a shear piezoelectric transducer that is pressed flush against one

edge of the sample with its axis of displacement aligned parallel to the length of the sample.

59
Care was taken to avoid destructive interference between echoes, which would result in

underestimating the amplitude of an isolated echo. This is a potential problem with the most

closely spaced echoes, which originate from the two pulses simultaneously emitted in opposite

directions by the meander coil. The problem is avoided by positioning the meander coil a quarter

wavelength away from a sample edge; then the two pulses essentially combine to form a single

pulse of approximately twice the duration of the input pulse to the meander coil, and no

destructive interference occurs.

During all types of measurements, the samples were mounted between C-shaped pole

pieces of an electromagnet (see Fig. 2.7.2). Samples were positioned carefully in the

electromagnet in such a way as to minimize pulling on the ends of the sample due to magnetic

attraction between the sample and the pole pieces. The amplitude of the SHq wave was

monitored as the magnetic field is decreased from high levels.

In Setup 2 shown in Fig. 2.7.2, the piezoelectric transducer was replaced by a second

EMAT. This setup was more suitable for measuring the wave amplitude at variable temperatures,

since the repeatibility of measurements with a piezoelectric transducer suffers from temperatureinduced shifts in the bond. The disadvantage of this setup is that it is more difficult to extract the

magnetostriction from measurements of the wave amplitude, since the wave amplitude is a

function of both the generation and reception efficiencies of the EMATs, which depend upon

magnetostriction in different ways.

Measurements of the ultrasonic wave amplitude i|r were performed on 1 22 x 2 1 x 9 mm

bars of A710 steel (samples CA, C5, and C7) and A533B (samples 1-2 and W-1). For each

sample, all measurements were performed on a single surface that was wet-ground using a sequence of sandpapers ranging in roughness from between 60 and 800 grit. Subsequently, at least 30 to 40 jum of material was removed by chemically polishing the surface with an aqueous

solution of oxalic acid and hydrogen peroxide. Between measurements, samples were stored

under a roughing pump vacuum, and, after the samples were exposed to air for a few hours during

the measurements, the chemical polishing of samples was repeated.

The magnetization, magnetostriction, and transverse incremental permeability were measured separately. The B-H curve of a sample was obtained by measuring the voltage induced

across a solenoid wound around a sample. The transverse incremental permeability was measured

through an eddy current technique suggested by Rose et al. [2.7.5]. These authors have shown

that the transverse incremental permeability iij is proportional to the unique frequency fg that

leaves the real inductance of a coil unaffected when the coil is brought close to a ferromagnetic

material. In the present work, a flat, spiral, oval shaped coil was used, and fg was measured as a magnetic field was applied. (The frequency fg varied between -10 kHz near saturation to -250

kHz at low fields.) The coil was oriented in such a way as to measure the incremental

permeability mostly along a direction perpendicular to the applied static field (that is, parallel to H„ in Fig. 2.7. 1 ). With this configuration, [^^ is expected to be B/H near saturation, and ^t is

60
derived as a function of the applied field by scaling fg so that [^^ equals B/H at high fields (with

values of B taken from the B-H curve and H expressed in units of A/m).

2.7.2 Measurement of A710 Steel

The purpose of tests on this material was to demonstrate that micromagnetic

measurements are sensitive to the formation of fine CRPs that strengthen the steel. Results are summarized in Fig. 2.7.3. The transverse incremental permeability (Fig. 2.7.3(a)) and the SHq

wave amplitudes (Fig. 2.7.3(b) and 2.7.3(c)) are different for the three samples. Differences

among the transverse incremental permeabilities occur primarily at low Hq, while differences

among the wave amplitudes extend to high fields. Figure 2.7.3(d) shows estimates of | A | obtained by substituting data from Figs 2.7.3(a) and 2.7.3(b) into eq (4), and, again, it is seen that

the curves are different for all three samples. The differences shown in Fig. 2.7.3(d) agree

qualitatively with results reported by other investigators [2.7.6]; for example, they observed that

the maximum static magnetostriction of copper-alloyed (1 mass percent) ferritic steel (WB 36)

reaches a peak value when fine CRPs form (similar to the case of sample C5).

Several aspects of these curves may be correlated with other measurements of the

samples. For example. Fig. 2.7.4 shows that the maximum transverse incremental permeability

correlates with the lattice strain determined from x-ray diffraction. Average values of the most

valid estimates of |A,| in Fig. 2.7.3(d) (those in the near-saturation regime, which lies above 25

kA/m in this case) also correlate with lattice strain. The same two quantities also correlate with

hardness (see Fig. 2.7.5). Figure 2.7.6 shows that the Hg-values of the maxima in Fig. 2.7.3(d)

correlate with the proportional limits of the samples.

The way that certain features of the magnetostriction and transverse incremental

permeability change monotonically with lattice strain and hardness suggests that the changes are linked to the population of strengthening CRPs. These micromagnetic measurements are therefore sensitive to the formation of these precipitates, which points to the possibility of using

either technique as an NDE tool for monitoring the evolution of CRT's in this particular steel.

2.7.3 Measurements on A533B Steel.

Micromagnetic measurements are performed on two specimens of this steel, one cut from

the base metal (1-2) and the other from a weld region (Wl). Results of room temperature

measurements are shown in Figs. 2.7.7 and 2.7.8. The transverse incremental permeability curves of the two samples practically overlap (the uncertainty of data points is ±5%), and the SHq wave

amplitudes largely agree with each other within estimated uncertainties. [It should be noted that, even though the receiver gain is increased for these measurements, much smaller wave amplitudes

are recorded in these samples than in the A710 samples (see Fig. 2.7.3(b)).] This close agreement of measurements in the two samples is consistent with the small difference of lattice strains

of these samples (e(l-2) = 0.00073 ± 0.00001, e(Wl) = 0.00077 ± 0.00001, and Ae/e = 5 ± 2%),

61
which reinforces the notion of micromagnetic quantities correlating with lattice strain, as suggested by measurements of the A710 samples.

These samples were also measured below room temperature in order to determine the

temperature dependences of the micromagnetic quantities and to ascertain whether they may be

used to predict the ductile-to-brittle transition temperature (DBTT) of this steel. The samples were cooled by placing them in spring-loaded contact with brass blocks carrying liquid nitrogen,

and nitrogen gas was sprayed on the samples to help suppress build-up of frost.

While the B-H curves of both samples do not change significantly between - 120 °C and

room temperature, both the transverse incremental permeabilities (Figs. 2.7.9 and 10) and the SHq

wave amplitudes (Fig. 2.7.1 1) of the samples do depend on temperature. The transverse

incremental permeabilities change with temperature primarily at low Hq (similar to effects of Cu

precipitation discussed previously). The temperature dependences of the maximum values of the

transverse incremental permeabilities are plotted in Fig. 2.7.12(b). The wave amplitude seems to be more sensitive to temperature; it displays a temperature dependence over a larger range of

fields Hq. Features of the most reliable data, portions of the curves lying in the near-saturation

regime (above ~25 kA/m), depend linearly upon temperature; for example, Fig. 2.7.12(c) shows

that the wave amplitudes at 80 kA/m decrease linearly with increasing temperature.

The temperature dependences of the two samples are different. Hence, even though at room temperature the samples share similar transverse incremental permeabilities and wave

amplitudes, one may measure these quantities at low temperatures to distinguish the samples. For

example, Fig. 2.7.8 shows that the wave amplitudes at - 128 °C are distinct. The temperature

dependences of the impact toughness of the two samples also are different, and the toughness

changes with temperature the most in sample 1-2, similar to the transverse incremental

permeability and wave amplitude (see Figs. 2.7.12(a) through (c)). The possibility arises, then,

that the temperature dependences of all three quantities are affected by common physical

properties. However, the fact that the micromagnetic curves do not reproduce the sigmoidal

shape of the impact toughness curve (that is, do not show evidence of a transition at the DBTT)

suggests that the relationship between the micromagnetic quantities and impact toughness is tenuous. Micromagnetic measurements cannot be used to determine directly the DBTT.

2.7.4 Conclusions

Micromagnetic measurements may be used to monitor the evolution of CRPs in A710

steel. Certain features of the magnetostriction and transverse incremental permeability are shown

to correlate linearly with the lattice strain, hardness, and proportional limit. The magnetostriction

is measured through use of EMAT-generated waves, a noncontacting technique which would be

necessary for determining zones of embrittlement in pressure vessels. Investigations in A533B

steel reveal that, although it is possible that the temperature dependences of micromagnetic

quantities and impact toughness depend upon common physical properties, micromagnetic

techniques cannot be used to determine directly the DBT temperature.

62
2.7.5 References

2.7.

Thompson, R.B., New configurations for the electromagnetic generation of SH waves in

ferromagnetic materials, in: 1978 Ultrasonics Symposium Proceedings, Cat. No.

78CH1344-1SU, J. deKlerk and B.R. McAvoy, eds., IEEE Press, New York (1978).

2.7.2 Thompson, R.B., Generation of horizontally polarized shear waves in ferromagnetic

materials using magnetostrictively coupled meander-coil electromagnetic transducers,

Appl. Phys. Lett. vol. 34, pp. 175-177, Jan. 1979.

2.7.3 Thompson, R.B., Physical principles of measurements with EMAT transducers, in: Ultrasonic Measurement Methods, Vol. 19 of Physical Acoustics, R.N. Thurston and

A.D. Pierce, eds.. Academic Press, Boston (1990), ch. 4, pp. 157-200.

2.7.4 Igarashi, B., Alers, G.A., Purtscher, P.T., An ultrasonic measurement of magnetostriction,

in: 1997 Ultrasonics Symposium Proceedings, IEEE Press, New York (in press).

2.7.5 Rose, J.H., Tai, C.-C, Moulder, J.C., Scaling relation for the inductance of a coil on a

ferromagnetic half-space, J. Appl. Phys., in press.

2.7.6 Altpeter, I., Becking, R., Jaschan, A., Kern, R., Lang, M., Rosner, H., Proof of Cu

Precipitates by Means of Micromagnetic Test Procedures, IZFP (Frauhofer-Institut fiir zerstorungsfreie Priifverfahren) Report No. 970137-TW, 1997 (unpublished).

63
H.^-' Steel sample Shear Horizontal Wave

(Top View)

Fig. 2.7.1 Setup for magnetostrictively generating SH ultrasonic waves with an EMAT. The meander

coil carries ac current I, and Ho is applied by an external magnet The resultant elastic displacements are

parallel to the sample surface.

Gated

Amplifier

Preamplifier

and Receiver

Setup 1

Gated

Amplifier

Preamplifier

and Receiver

N

Setup 2

Fig. 2.7.2 Experimental setups for measuring the amplitude of SHo ultrasonic waves.

64
(a) Transverse Incremental Permeability

Ho(kA/m)

Fig 2 7 3 (a) Transverse incremental permeabilities, (b) SHo wave amplitudes measured m Setup

1 of Fig 2 7 2 (c) SHo wave amplitudes measured in Setup 2 of Fig. 2.7.2, (d) estimates of j A | from eq'(4) and (e) B-H curves obtained in the CA(0), C5(a), and C7(n) A710 samples.

65
0.00045 0.00055 0.00065

-r240

--200

-- 160

120

0.00075

>

X

CO

Lattice Strain

Fig. 2.7.4 Estimated \X\ averaged over high Ho(0) and maximum transverse incremental

permeability (a) versus lattice strain.

52 54 56

l-lardness (RA)

58 60

Fig. 2.7.5 Estimated | X \ averaged over high Ho(0) and maximum transverse incremental

permeability (a) versus hardness.

66
20 T ^

1 50 200 250 300 350 400

Prop. Limit (l\/IPa)

450 500

Fig. 2.7.6 Ho at maximum |X| in Fig. 2.7.3(d) - (0) versus proportional limit.

180

160 5

"k RoomTenperature Measurement

140- \ ~^^'^

•^ 120-

•^ 100-

2 80-

i 60- ^^^^^-oa^

°^°^^°^^^—

40- -

20-

0- 1 1 —1

1 1 1 1

1 10 20 30 40 50

Ho (kA/m)

60 70 80

Fig. 2.7.7 Comparison of tranverse incremental permeabilities of A533B samples 1-2 (o) and

Wl(n) at room temperature.

67
40 T

SHo Wave Amplitude at Room Temperature

200 T

10 20 30 40 50

Ho (kA/m)

60 70 80

Fig. 2.7.8 Comparisons of SHq wave amplitudes measured in A533B samples 1-2 (o) and Wl ()

at room temperature (top) and at - 128 °C (bottom).

68
A533B (1-2)

-0-24 "C

-D--20 "C

-6- -77 ''C -X--120 °C

"^2^1

'"Kqv.>

^^h-U^s,

^^41-R-a^.x•^2«»i?lw:^-^<xl

30 40 50

Ho (kA/m)

60 70 80

Fig. 2.7.9 Transverse incremental permeability of sample 1-2 versus Hq at 120 °C (x), 77 °C

(a), -20°C(n),and24°C(o).

A533B(W1)

10 20 30 40 50

Ho (kA/m)

60 70 80

Fig. 2.7. 10 Transverse incremental permeability of sample Wl versus Ho at - 1 16 °C (), 60 °C

(a), and 25 °C (o).

69
250 T

10 20 30 40

Ho (kA/m)

50

10 20 30 40 50

Ho (kA/m)

60

60

70 80

128 °C

70 80

Fig. 2.7.1 1 SHo wave amplitude measured at various temperatures in sample 1-2 (top) and

sample Wl (bottom) using Setup 2 in Fig. 2.7.2.

70
250t (a) Impact Toughness

100

>

180 T

170

O^

160 -

atX

CB 150 +

140

-ISO -100 -50

(b) Maximum Transverse

Incremental Permeability

50 100

250t

-150 -100

(c) SH Wave Amplitude

-SO Temperature {^C)

50 100

Fig. 2.7.12 Comparisons of temperature dependences of impact toughnesses (a), maximum

transverse incremental permeabilities (b), and SHq wave amplitudes at 80 kA/m (c) measured in

A533B samples 1-2 (o) and Wl (D).

71
2.8 DISLOCATION MOBILITY UNDER A STATIC STRESS—Ward Johnson

2.8.1 Introduction

Ultrasonic measurements can be highly sensitive to the density, symmetry, and pinning of

dislocations in crystals because the nonelastic response of dislocations to an acoustic wave causes damping and slowing of the wave. This phenomenon was explored in this project as a potential

tool for nondestructively evaluating the mechanical properties of irradiated steel, especially the

ductile-to-brittle transition temperature. The principal experimental approach employed was to use resonant ultrasonics to sense the degree to which dislocations break away from pinning points

when surrogate materials were subjected to loads in a tensile testing machine. This type of

measurement provides information about the mobility of dislocations.

2.8.2 Experimental Apparatus

A unique system was implemented for measuring resonant ultrasonic properties of

specimens undergoing static or dynamic loading in a tensile testing machine. Figure 2.8.1 shows

the configuration of the specimen and the grips that are mounted in the testing machine.

Elecromagnets around the machine grips provide a static magnetic field along the specimen axis.

Liquid nitrogen cooling and resistive heating of the grips provide temperature control from

- 100 °C to +70 °C. Highly accurate damping and velocity measurements are achieved with a

cylindrical "trapped-mode" geometry that has resonant vibrational modes localized in a central

portion of the specimen and, thus, has insignificant energy losses through the testing machine

grips [2.8.1]. This geometry is shown in Fig. 2.8.2. The length of the central trapping region / is 51.7 mm for A710 (C5 and C7) specimens and 57.5 mm in the A533B (1-1, 1-2, and W-1)

specimens. The total length Z is 157 mm for A710 and 166 mm for A533B. The losses of energy

through transducer coupling are essentially eliminated by employing noncontacting

electromagnetic-acoustic transducers [2.8.2]. The sample is driven into resonant vibration by a long tone burst, and the frequency and amplitude of vibrations are measured during the

subsequent ring-down. A detailed description of this technique of signal analysis is presented

elsewhere [2.8.3]. Because of the high accuracy and the frequencies employed (typically, 0.8 to 3

MHz), this system is very sensitivity to the anelastic properties of dislocations and associated

microstructural changes.

Radial and flexural resonant modes show similar changes in damping in response to static

loading, but have different values for the equilibrium background damping. In this report, the

mode types and frequencies of the measurements will be noted only in the figure captions.

The field provided by the electromagnets is employed in acoustic transduction. It also

reduces damping associated with magnetoelastic responses within the material, making the

measurements more reproducible and sensitive to dislocation effects. Since dislocations were the

principle focus of this part of the feasibility project, results relating to magnetic damping are not

72
presented here. All of the measurements in this report were performed with the magnetic field in

the range 28 to 32 kA/m (350 to 400 Oe).

2.8.3 Results

a. Before plastic deformation. Figure 2.8.3 shows the damping at several ultrasonic

frequencies in three A533B specimens prior to loading. As described elsewhere, the 1-1 and 1-2

specimens were taken from different sections of base metal, and the W-1 specimen was taken

from weld metal. The damping is expressed as the dimensionless logarithmic decrement 6, which

is the reciprocal of the number of cycles for the free-decay amplitude of resonant vibrations to drop to 1/e of its initial value. The damping in all three specimens increases monotonically with

frequency. This is consistent with the frequency dependence expected for dislocation damping in

impure materials at these frequencies [2.8.4]. The data suggest that the density of dislocations

that move in response to acoustic oscillations is lowest in 1-1 and highest in W-1.

A correlation between the ultrasonic results and Charpy results is pursued in Fig. 2.8.4 by

taking the slopes of 6 versus frequency for each specimen, obtained from the linear regressions

shown in Fig. 2.8.3, and plotting these versus the temperature at which the Charpy impact energy

drops to one half that of the upper plateau. The fact that 1-1 becomes brittle at higher

temperatures than 1-2 and W-1 is consistent with the lower mobile dislocation density of 1-1 suggested by the ultrasonic damping measurements.

As discussed in section 2.8.4, accurate absolute measurements of damping in steel would be

very difficult, if not impossible, on an actual pressure vessel. Therefore, in further pursuit of a

practical nondestructive ultrasonic technique, a series of tests were developed to determine the

relative changes in damping that occurs when specimens are subjected to tensile loading.

Ultrasonic damping and velocity were measured as the load was ( 1 ) stepped rapidly from near

zero to a significant fraction of yield, (2) held at the higher load for periods ranging from 10 to 600 s, and (3) reduced back to the original low load and held there. If dislocations are broken

away from pinning points by the application of a load, one expects the damping to increase

immediately but to recover with time as thermally activated repinning takes place. This effect is observed in materials that have been previously plastically deformed (see below). However, no

effects of this type were observed in undeformed A533B (1-1, 1-2, and W-1) or A710 (C5) at temperatures ranging from -80 to 70 °C.

b. During plastic deformation. Ultrasonic measurements of two specimens of A710 steel, C5

and C7, were performed as the load was increased through yield. The loads were increased in

discrete steps, and, at each load, 6 was measured as a function of time. Figure 2.8.5 shows 6 at two times following the step to each value of the load. In C5 (Fig. 2.8.5(a)), 6 was relatively

constant as the load was increased until reaching approximately 450 MPa, where time dependent

changes occurred after the load step. Increases in 6 after the steps were greater at higher loads,

but always approached a constant with extended time at a given load. The changes in 6 are interpreted as being the result of the depinning of dislocations and the increase in dislocation

73
densities that accompany plastic deformation. Similar effects have been reported by numerous

researchers for a variety of metals and nonmetals. The continuous nature of the changes shown

for C5, and the loads where they occur, are similar to the plastic strain shown in the conventional

stress-strain curve obtained for this continuously yielding material. (See data presented in Fig.

2.2.1.) Measurements of the diameter of C5 before and after this test showed that a plastic strain

of 0.15% had been introduced. Sample C7, which was found to yield discontinously in the

conventional stress-strain test, had a relatively abrupt increase in 6 near 495 MPa (Fig. 2.8.5(b)).

This load is somewhat higher than the 47 1 MPa yield stress found in the conventional tensile test, probably because the ultrasonic test was performed much more slowly. Unlike C5, changes in 6

proceeded more rapidly as the load was held at the highest load, consistent with discontinous

yielding.

c. After plastic deformation. In plastically deformed A533B and A710, reproducible

changes in 6 are observed with loads well below those where macroscopic deformation had

occurred in previous tests. Figure 2.8.6 shows measurements on a specimen of material 1-1 before and after 0.39% plastic prestrain. For the prestrained material, the load was initially held

near 6 MPa for 1000 s, then increased in 0.5 s to 163 MPa (35% of the yield stress, as determined

from the stress-strain test), held at that load for 1 00 s (during which time no ultrasonic

measurements were performed), and, then, decreased back to 6 MPa. The 100 s loading induced an increase in 6 followed by a slow recovery to the original value. A similar loading of material

with no plastic prestrain induced no load-induced changes in 6 even though the load was much

higher (444 MPa). The load-induced changes in the deformed material are interpreted as being a

result of breakaway and, perhaps, multiplication of dislocations followed, after release of the load,

by recovery involving thermally activated movement of dislocations and/or point defects. This

effect is found to be independent of the ultrasonic frequency. The solid line in Fig. 2.8.6 is a fit to the form 6 = 6^^+ [A-Bln(l + t/x)Y, which accurately matches changes in 6 that accompany

high-temperature isothermal recovery of cold-worked aluminum [2.8.5]. 6^^ is the equilibrium

decrement at 6 MPa before the 1 63 MPa load. The success of this function in fitting the data of

Fig. 2.8.6 suggests a mechanism similar to that of high-temperature recovery, involving

dislocation annihilation followed by movement into a lower-energy configuration.

The magnitude of load-induced changes in 6 increases with the magnitude of the load, as one would expect from dislocation breakaway. Figure 2.8.7 shows this dependence in a specimen of

C5 that was previously deformed 0.15%. The response to a load also depends strongly on

temperature. Measurements were performed on 0.15%-deformed C5 at several temperatures

from -90 °C to +27 °C before and after loading at 420 MPa for 10 s. Figure 2.8.8 presents these

measurements with the equilibrium 6^^ subtracted. At the higher temperatures, the initial change

in 6 is higher and the recovery is most rapid. This behavior is consistent with a mechanism

involving thermally assisted unpinning of dislocations. That is, at the higher temperatures,

dislocations are more rapidly unpinned during the period at 420 MPa, and, after the loading,

thermally activated recovery proceeds more rapidly.

74
Since fracture behavior is also determined by the kinetics of dislocation unpinning, one may

expect a close correlation between such ultrasonic results and the ductile-to-brittle transition.

2.8.4 Conclusions

The results presented here indicate that ultrasonic damping may provide information about

mechanical properties by sensing the anelastic behavior of dislocations in equilibrium and under

static loading. The equilibrium 6 of undeformed A553B shows a correlation with the

ductile-to-brittle transition in the three specimens studied (Fig. 2.8.4), and the sign of the

correlation is consistent with that expected for dislocation effects. Increases in 6 during plastic

deformation provide information about yielding similar to that of a stress-strain curve. In

plastically deformed material, the recoverable temperature-dependent ultrasonic response to loading is, apparently, sensitive to the kinetics of dislocation unpinning and, thus, may be closely

related to the ductile-to-brittle transition.

From a practical perspective, however, the results are less than encouraging. Although

interesting from the perspective of materials characterization, the techniques employed during and

after deformation are useless for nondestructive evaluation because they involve modification of

the material. The measurement of equilibrium 6 before deformation is nondestructive. However,

the values obtained for the equilibrium 6 (Fig. 2.8.3) are so low that they are probably impossible

to measure in the field. Because of the extensive complicated geometry of a pressure vessel and

its structural supports, a high-resolution resonance method can not be employed. On the other

hand, conventional pulse-echo methods suffer from diffraction losses that, even in plates with

polished parallel surfaces, typically result in values of 6 two orders of magnitude larger than the

intrinsic 6 measured here. Considering that the surfaces of pressure vessels (especially the

stainless cladding) are highly irregular, scattering losses would be even more severe and small

changes in intrinsic damping would be undetectable. Perhaps, other more exotic approaches to measuring damping would be useful to pursue. However, at present, there is no known technique

that can measure damping comparable to that shown in Fig. 2.8.3 on an extensive irregular object

such as a reactor pressure vessel.

2.8.5 References

2.8.1 Johnson, Ward; Auld, B.A. and Segal, E., Trapped torsional modes in solid cylinders, J. Acoust. Soc. Am., Vol. 100, p. 285 (1996).

2.8.2 Johnson, Ward; Alers, George, Force measurement using vibrational spectroscopy,

Review of Scientific Instruments, Vol. 68, p. 102 (1997).

2.8.3 Johnson, Ward; Ultrasonic resonance of metallic spheres at elevated temperatures, J. Phys.

IV (France), supplement to J. Phys. Ill (France), Vol. 6, p. C8-849 (1996).

75
2.8.4 Granato, A.V.; Lucke, K.; Physical Acoustics, Vol. IVA, edited by W. P. Mason

(Academic, New York, 1966), p. 225.

2.8.5 Johnson, Ward; Ultrasonic damping in pure aluminum at elevated temperatures, J. Appl.

Phys., to be published (March, 1998).

76
Electromeignet

Electromagnet

2.8.1 Configuration of the specimen and grips mounted in a tensile testing machine. Each of the

grips is encircled by a heater, a liquid nitrogen cooling coil, and an electromagnet.

77
6.35rain radius 6.35niin radius

3/4"-12 thread

.lO.SOOmra 10.438inin .10.300mm

3/4"- 12 thread

19mnn 19mjn

2.8.2 Geometry of a cylindrical ultrasonic resonant trapped-mode specimen.

'O

2.5E-4

2.0E-4

1.5E-4

l.OE-4

5.0E-5

O.OEO

W-l weld

1-1 base metal

1 2

Frequency (MHz)

2.8.3 Equilibrium 6 versus frequency of radial resonant modes in A533B specimens before

deformation. 30 °C. 28 to 32 kA/m. The lines are linear regressions for each specimen.

78
>

o

t

CO

-30 -20 -10

Temperature at 1/2 maximum absorbed energy (**C)

2.8.4 Slopes of the linear regressions shown in Fig. 2.8.3 versus the temperature at which the

absorbed energy in the Charpy tests drops to one half that of the high-temperature plateau.

3.0

2.9

2.8 h 2.7 ^ 2.6 oC 2.5 o 2.4

2.3 22

2.1

2.0

IS 9 at load 300 sat load

,0"

^'

200 300 400 500

Stress (MPa)

(a)

600

J.ir

2.9

'

• ISs at bad

2.8 o 300 sat load •

2.7 . a 400 3 at bad a ^26 . oO 2.5 o O 2.4

* •• •

•

2.3

2.2

• • • • • • • •••-•SSS.

2.1 1 n200 300 400 500

Stress (MPa)

(b)

600

2.8.5 6 of the 0.82 MHz flexural mode of specimens C5 (a) and C7 (b) versus load during plastic

deformation (33 °C, 37.8 kA/m).

79
4E-4

3E-4

•o

2E-4

lE-4

o Before plastic defonnadon

A After 0.39% plastic strain

, jjif/^^^jij>/^

163 MPa for 100 s

444 MPa for 100 s

^UflaaAAAf

I o o o o ooo

2000 4000

Time (s)

6000

2.8.6 6 of the 1.55 MHz radial mode of specimen 1-1 (A533B) versus time before and after

0.39% plastic prestrain. The load was held at 6 MPa for 1000 s, increased to an elevated value

(444 MPa before and 1 63 MPa after plastic deformation), held at the elevated load for 1 00 s, then

decreased back to 6 MPa. No ultrasonic measurements were recorded while holding at the

elevated loads. The solid line is a fit of the recovery data to 6

28 to 32 kA/m).

1.6E-4

6, +[A- Bln(l+t/T)f. (30 °C,

1.5E-4

•o 1.4E-4

42000 46000 50000 54000

Clock time (s)

58000

2.8.7 6 of the 1 .55 MHz radial mode of specimen C5 versus time during a loading sequence. The

load was held at the elevated values indicated for 10 s, and held at 6 MPa at all other times. The

specimen was plastically prestrained 0.15% before the measurements were taken. No ultrasonic

measurements were recorded while holding at the elevated load. (30°C, 27.2 kA/m)

80
1.8

1.6

1.4

1 1.2

O

1.0

0.8

1 0.6

0.4

0.2

0.0

-I 1 1 I I r—

I

1

I "1 1 1 1 1 \

1 I 27 °C

^ -20 C

1000 2000 3000 4000 5000

Time since 420 MPa load (s)

2.8.8 Recovery of 6 of specimen C5 following 10 s/420 MPa loading at 27, -20, and -90 °C.

The equilibrium 6^^ (measured before loading at each temperature) is subtracted. (2.12 MHz

radial mode, 3 1 .0 kA/m)

81
2.9 MEASUREMENTS UNDER CLADDING

2.9.1 Magnetic Field Penetration of SNUPPS Cladding — G.A. Alers

In order to measure the magnetostriction coefficient of the A533B steel under the cladding,

it is necessary to subject the steel to an alternating magnetic field and an adjustable static bias

field. Unfortunately, the cladding is a conductor and is slightly magnetic so it can effectively

shield alternating magnetic fields. It will have no shielding effect on the static bias field except for

the physical separation it will impose between the source of the bias and the A533B surface. In order to estimate the effectiveness of the cladding as a dynamic current shield and to establish

quantitative design parameters for a transducer to measure the magnetostrictive characteristics of

the A533B pressure vessel material under the cladding, we consider the electromagnet

configuration shown in Fig. 2.9.1. A Hall probe capable of measuring dynamic magnetic fields

was placed either above or below the cladding sample midway between the pole pieces of the

electromagnet. It was oriented to measure the component of the field parallel to the surface of

the slab. If the electromagnet is driven by an alternating current at frequency f, we can define the

following measurables:

h(f) == the dynamic magnetic field under the cladding

T = the thickness of the cladding

f = the frequency of the electromagnet current and dynamic field

\{f) = the magnitude of the field on top of the slab and

H = a static magnetic field supplied by a distant magnet (not shown) to subject the region

under the slab to a bias field.

For a large spacing between the pole pieces of the electromagnet, classical electromagnetic

theory would predict

h(f) = h„ exp(-T/d) (1)

where d = the electromagnetic skin depth = k\/(p/|Lif). The constant k has the value 0.052 mm if the electrical resistivity p of the slab is measured in fiQ-cm and fin megahertz. The relative

permeability of the slab is |li which may not equal 1 for the slightly magnetic stainless steel cladding.

For a first-order estimate, the Metals Handbook gives (for type 308 stainless steel) p = 70

P-Q'cm, |i = 1. Equation (1) then gives

d=1.4mm@f=100kHz

d = 4.3mm@f= 10 kHz

d=14mm @f=lkHz.

82
or, for T = 10 mm, the dynamic field under the cladding will be attenuated to 37% of its value at the surface if the frequency is 2 kHz.

In order to get a more realistic measure of the field that will penetrate the cladding of a

pressure vessel, an experimental version of Fig. 2.9.1 was set up. The electromagnet had a laminated core and was driven by a power amplifier capable of delivering up to 14A (rms) at any

frequency from to 50 kHz into a properly matched load. A commercial Hall probe was

mounted in a wooden base so that it would measure the tangential component of the field at a

point midway between the C-shaped pole pieces. A 4 mm thick slab of the stainless steel cladding

cut from a sample of the SNUPPS pressure vessel was placed between the electromagnet and the

Hall probe. With this configuration, the magnetic field that penetrated the cladding could be

measured directly as a function of the current supplied to the electromagnet at any desired

frequency. Figure 2.9.2 shows the magnitude of the field observed under the slab at a frequency

of 0.1 kHz (100 Hz) for various currents in the electromagnet. The graph also shows the field

values observed for 5 mm thick slabs of other materials including air. As expected, the field is a

linear function of the current in the electromagnet and become smaller for a fixed current as the

conductivity of the slab increases from stainless steel to aluminum.

In order to estimate the frequency at which it will be practical to operate an alternating

current electromagnet above the cladding and still have a useful amount of magnetic field

penetrate to the A533B substrate, field values were measured under various metal slabs as a

function of the frequency of the current in the electromagnet. Such a measurement produces the

quantity h(f) in eq (1). The quantity h^, in eq (1) was measured separately with the Hall probe

above the metal slab as shown in Fig. 2.9.2. For a constant current in the electromagnet, h<, was

observed to increase with frequency because eddy currents induced in the slab under the pole

pieces produce a field that adds to the tangential field midway between the electromagnet pole

pieces while they subtract from the tangential field below the slab. Equation (1) predicts that a graph of ln(h/ho) versus the square root of the frequency should be a straight line with a slope of

1 19/(11/9) ^^^ ^^ intercept of zero. Such a graph is shown in Fig. 2.9.3 for measurements on 5 mm slabs of aluminum and lead as well as the 4 mm thick slab of cladding cut from the SNUPPS

RPV. The values of the slopes shown on the graph are for "best fits" to the measurements. The

table below compares the slopes with the predictions of eq (1) using values of the resistivity of

aluminum, lead, and type 308 stainless steel taken from handbooks.

Since the calculated and measured slopes are in reasonable agreement, eq (1) can be used to estimate the attenuation of the magnetic field to be expected as a function of frequency for

cladding materials of various thicknesses. In particular, the solid line in Fig. 2.9.3 describes the

attenuation for a realistic thickness of cladding of 8 mm. Based on this solid line, we estimate that

8 mm of cladding would attenuate the magnetic field to 37% (1/e) of its value at the surface if the

operating frequency were 4 kHz or to 20% of the surface value if the frequency were 10 kHz.

83
Table 2.9.1 Comparison of measurements with the predictions of Eqn (1)

Material

Thickness

(mm)

Slope

(Fig. 2.9.3)

Calculated

Slope

Resistivity

(|iQ-cm)

Permeability

(Relative)

Aluminum 5 -65 -59 2.8 1

Lead 5 -23 -21 20 1

Cladding 4 -8 -9 70 1

For a cladding thickness of 8 mm and accepting a factor-of-10 attenuation in the dynamic

field by the cladding, Fig. 2.9.3 indicates that an acoustic wave at a frequency of 10 kHz could be

excited by magnetostrictive coupling in the A533B alloy under the cladding from an AC

electromagnet mounted on the cladding. The acoustic wave would be a shear wave with a propagation velocity of 3 mm/jis so its wavelength at 10 kHz would be 300 mm. This is approximately twice the thickness of the pressure vessel wall so establishing a standing wave

resonance in the thickness dimension would be quite possible. A measurement of the amplitude of

this resonant vibration at a fixed current in the electromagnet as a function of a biasing field would

constitute a measurement of the magnetostrictive efficiency as a function of magnetic bias which

is the information needed to obtain the parameters | A, | at high Hq and Hq @ max | X \ required for

predicting the strength properties of interest (see Figs. 2.7.5 and 2.7.6).

2.9.2 Static Magnetic Properties (B-H Looper) — F.R. Fickett

The electromagnet and Hall probe configuration shown in Fig. 2.9.1 could be used with a

direct current in the electromagnet to deliver a static magnetic field to the A533B steel under the

cladding. Its magnitude and distribution could be calculated by finite element techniques if sufficient information about the geometry and the cladding properties were available. Such a

calculation would use measurements of the magnetic field under the pole pieces and in the space

between the pole pieces at the air-to-cladding interface to deduce a hysteresis curve for the

A533B steel. In order to determine if such measurements would be possible on an actual pressure

vessel, the Hall probe and electromagnet shown in Fig. 2.9.1 were placed on the clad surface of a

1 m square section of the Shoreham RPV and the fields around the electromagnet were measured

as a function of the current in the electromagnet. The results showed that a measurable hysteresis

from the A533B base metal could be detected through the cladding. Thus, it is possible to measure the coersive field He of the A533B steel under the cladding although the accuracy

obtainable could not be judged.

84
2.9.3 Nonlinear Ultrasonic Experiments for Measuring RPV Wall Properties —D. Hurley

Nonlinear ultrasonic experiments indicate that the third-order elastic constants of steels with

copper-rich precipitates are sensitive to microstructural changes that occur during hardening. One

measure of the third-order elastic properties is the nonlinear ultrasonic parameter p. In the

laboratory experiments described in Sect. 2.5, p correlated with hardness in a series of steel samples. Due to geometrical considerations, it may be more practical to measure the third-order

properties in the reactor itself by a means other than p. One possibility is to perform ultrasonic

mixing experiments such as described below. Preliminary estimates indicate that this approach

may be feasible using ultrasonic waves in the 5 to 10 MHz range. However, further calculations

and additional laboratory experiments are necessary.

Figure 2.9.4 depicts the basic concept of a mixing experiment. Two transversely polarized

ultrasonic waves at frequencies f, and fj are generated at different positions on the clad inner

surface of the reactor. The waves penetrate the stainless steel layer (refraction has been ignored

for simplicity), reflect off the reactor outer wall, and propagate back towards the reactor inner

surface. The two waves will intersect inside the reactor steel at a position which depends on the

incident angles and transducer separation. At the intersection, a longitudinal wave of frequency f, + fj will be generated by the same nonlinearities of the interatomic forces that give rise to the

coefficient p. The longitudinal wave thus generated propagates at an angle determined by the

sum of the transverse wavevectors. As shown in the figure, the transverse-wave transducers are arranged so that the longitudinal wave propagates back to the reactor's inner surface where it could be detected by a third transducer.

The amplitude of the generated longitudinal wave depends on the third-order elastic

properties of the intersection volume. However, the amplitude also depends on other experimental

parameters, such as the transducer dimensions, the angle between beams, the incident wave

amplitude and frequency, and the second-order elastic constants. Therefore, thought must be

given to an experimental approach which will isolate the effect of the third-order elastic properties

from other effects. For instance, the behavior of amplitude as a function of incident angle or frequency may provide useful information. By varying the distance between the two generating

transducers, the properties as a function of depth can be investigated. The use of mixing

techniques (as opposed to harmonic generation) ensures that the detected signals are due solely to the third-order properties of the interaction volume and are not related to experimental artifacts.

85
Electromagnet

Fig. 2.9.1 Schematic diagram of an electromagnet designed to apply an alternating magnetic field

to the cladding on the inner surface of a reactor pressure vessel. The field h(f) penetrates the

cladding and is applied to the pressure vessel itself

Field vs Current @ 100 Hz

ivv mm Air=ho(f) v ^^

90 - J>^ ^mmSS

80 ^y^ f^ , "5 mm Air

f

70

60 ^X^-' m * ^x

o

0) iZ1

50

40

y\ * ^* ^, ^ b mm Al

S 30

20 ^^ *

10 '^- *x

01^ ^ 1. 1 ) H ^—

^

u- ^ 1 \

1

'

1 3 4 5 6 7 8 Magnet Current @ 100 Hz (rms Amps)

10 11

Fig. 2.9.2 Magnetic field that penetrates slabs of different materials as a function of the current in

the electromagnet shown in Fig. 2.9.1.

86
-0.5

-1

-1.5

-2.5

^=t:—

Field Attenuation

[Frequency in MHz^.S

(I Vi>0T - »^2. ^0.03 0.04 ^ 0.05 0.06 0.07 0.08 0.09

\

Slope = -65 \

\

Slope = -23

\

^ \

8mmSS

\

\

\

A.

0.1 Frequency (kHz)

10

Fig. 2.9.3 Magnetic field penetrating slabs of different materials as a function of the frequency of

the 1 A (rms) current in the electromagnet. The solid line depicts the field expected for a cladding

thickness of 8 mm. (a Aluminum, Lead, Stainless Steel)

Stainless

cladding

reactor

steel

Fig. 2.9.4 Schematic diagram of a proposed configuration of ultrasonic transducers that could measure the nonlinear acoustic response of RPV material under the cladding by measuring the

efficiency of mixing of two input waves to produce a third wave.

87
3. DISCUSSION

3.1 Physical Property-Strength Correlations

A major subtask of the research presented here was to produce a collection of physical

property measurements on a material of controlled microstructure for which the relationship

between microstructure and mechanical strength was known. The A710 (Fe-1 .13 mass % Cu

alloy) provided a steel whose strength properties could be varied over an extensive range by heat

treatments and whose changes in microstructure were similar to those thought to be the cause of

radiation induced embrittlement of A533B steels. Therefore, the A710 steel was used here as a

surrogate of A533B on which to establish relationships between destructively measured strength

properties and nondestructively measured physical properties. Tables 3.1.2, 3.1.3 and 3.1.4 at the

end of this section compile the results of all of the mechanical strength measurements and the

physical property measurements described in the preceding sections of this report. Examination

of all the data shows that there were 6 strength-related properties and 1 physical properties

available for establishing interrelationships. These are listed in Table 3.1.1 along with the symbols

used to describe them in the text.

In performing the mechanical strength tests on the A710 surrogate alloy as well as on the A533B

RPV material, we observed a major difference in deformation mode and its influence on the

development of relationships between physical properties and strength must be kept in mind.

Specifically, the solution treated and peak-aged samples of the A710 alloy showed a smooth

transition from elastic to plastic strain as the stress was increased. On the other hand, the over- aged A710 steel and the A533B samples showed discontinuous yielding with little or no

Table 3.1.1 List of mechanical strength properties and physical properties measured in this

program.

Strength property Symbol Physical property Symbol

Proportional limit PL Elastic modulus E, G, or B

0.2% yield strength 0.2%YS Density P

Ultimate Tensile Strength UTS Internal friction Q-' or 6

Ductile-to-brittle Nonlinear ultrasonic parameter P

transition temperature DBTT Internal strain e Hardness HRA High field average magnetostriction X°°

Elongation % Elong. Field at maximum magnetostriction TJ Max. transverse incremental permeability ^, (max)

Magnetic coercive force He

Saturation magnetization ^^pM,

88
plastic flow preceding the load drop when loaded at a constant strain rate. All the materials

showed similar ultimate tensile strengths and large elongations so they all had comparable

ductilities. Thus, a difference in correlations with physical properties can be expected for strength

parameters that reflect yielding phenomena at low plastic strains, and ductile and work-hardening

parameters that describe deformation at high plastic strains. Figure 3.1.1 graphs three important

strength parameters plotted against hardness for the A710 steel. It shows a good linear

correlation or proportionality between the hardness and the UTS but a nonlinear relationship with

the proportional limit and the 0.2%YS. Plotting the DBTT versus hardness for the A710 alloy

series shows a linear correlation similar to that observed for the UTS in Fig. 3.1.1. Thus, if a nondestructive measurement of the DBTT is desired, then those physical properties that show a

strong correlation with hardness or UTS can be considered good candidates. If a physical

property does not correlate well with hardness, it may be useful as a predictor of the proportional

limit or the 0.2% YS and reflect the early stages of plastic flow when dislocations are just

beginning to be mobile.

To test for strong correlations between physical properties and the strengthening effect of

the copper rich precipitates (CRTs), each physical property was plotted against the hardness and

against the proportional limit for the A710 alloy in its three conditions of tempering. Since the

physical properties have a wide range of numerical values, these plots are best presented to the

reader in a normalized format in which the values are divided by the value for the solution treated,

initial condition of the alloy. This form of presentation allows the percentage change in property

value to be read from the ordinate as the hardness or the proportional limit (the abscissa) extends

from its minimum to its maximum value. Figure 3.1 .2 shows the case of plotting against the

hardness and Fig. 3.1.3, the case for plotting against the proportional limit. The choice of which

property was plotted against which strength parameter was motivated by the requirement that the

resulting curve need not be linear but should be monotonically increasing or decreasing in order to avoid double valued relationships. The fact that the correlations with hardness shown in Fig.

3.1.2 are all increasing while all those correlated with the proportional limit (Fig. 3.1.3) are decreasing is coincidental. Examination of these graphs shows that 8 of the 10 properties

measured and listed in Table 3.1.1 showed a monotonic correlation with either the hardness or the

proportional limit. The two properties not graphed, the density and the saturation magnetization,

showed very small changes during the heat treatment and would be difficult to measure in a

nondestructive test.

Figures 3.1.1 and 3.1.2 show hardness as the abscissa. Because it is equivalent to the

ductile-to-brittle transition temperature, the figures demonstrate that there are four physical

properties that could be candidates for a nondestructive predictor of the DBTT. These are the

internal strain e, the high field average magnetostriction X„, the maximum transverse magnetic

permeability ^t (max) , and the nonlinearity parameter p. All of these, with the exception of e and

jif, can be measured by ultrasonic techniques. As discussed in Sect. 2.9, it is also possible to make

these measurements through the stainless steel cladding so they could, in principle, be performed on the inner surface of a pressure vessel. To emphasize this encouraging result. Fig. 3.1.4 graphs

these three parameters against the DBTT. Here the normalization factor used is the value of the

89
property for the alloy in the overaged condition (C7). These results describe only three heat

treatment conditions of a steel containing 1.13 mass percent Cu in which the responses to the

Charpy test were dramatically different (see Fig. 2.2.3). In RPV steels, the copper

concentrations are much smaller and the shelf energies and transition temperatures are much

closer together. However, the properties measured appear sensitive to the strength parameter in

that they change by more than 1 0% over the range of hardness and DBTT investigated.

The graph of properties against the proportional limit (Fig. 3.1.3) show that the shear

modulus and the coercive force, H^, are insensitive to the heat treatment that forms the CRPs and

maximizes the hardness. Most of the change in these properties occurs when the copper forms e phase precipitates. The internal friction 6 and the field at which the magnetostriction is maximum

(H^ax) show a good correlation with the proportional limit and sensitively indicate the change in

this strength parameter during the formation of the CRPs.

3.2 Dislocation Mechanisms

Concerning the dislocation mechanisms that may underlie these correlations, the mechanical

properties involving large deformations (hardness, UTS, and DBTT) are correlated with physical

properties (e, p, [i^, and X°°) that reflect changes in microstructure on a nearly atomic scale. The

internal strain, e, is deduced from x-ray diffraction line broadening and is interpreted as reflecting

the strain around the CRP clusters as they retain coherency with the lattice. This internal strain

field obviously interacts with the dislocations to impair their motion over long distances as measured by the UTS. The acoustic nonlinearity parameter P measures the changes in the

interatomic forces produced by strain. Hence, it is not surprising that it, too, is sensitive to internal strain fields and, thus, can be used as a nondestructive measure of internal strains. The

two magnetic parameters, |n, and X°°, measure the mobility of domain walls and the freedom of

atomic spins to rotate relative to the direction of the applied field, respectively. They, too, are known to be sensitive to applied strains and, therefore, their measurement can be used to measure

internal strains. The fact that all four properties are related to the dislocation motion at the

ultimate tensile stress indicates that internal strains control the motion of dislocations at high

levels of deformation.

For the low strain mechanical properties (proportional limit and upper yield stress), the

dominant mechanism of dislocation motion is the release of dislocations from pinning points or impurity atmospheres by an applied stress. Since H^^^ ^^d 6 change with the proportional limit

during the formation of CRPs as shown in Fig. 3.1.3, these two measurable parameters must be

sensitive to the depirming of dislocations and the initiation of plastic flow at small strain levels.

The sensitivity of 6 to the release of dislocations was measured directly in our measurements of

the effect of a static stress on dislocation mobility that are described in Section 2.8 above.

Unfortunately, these experiments show that a measurement of the static stress needed to increase

the internal friction or to produce a small amount of plastic flow cannot be considered a nondestructive test because they leave the material changed and require application of very high

stresses to activate dislocation motion. On the other hand, the time dependent recovery of 6

90
following microscopic plastic flow shows some promise of indicating the degree of dislocation

mobility that is available at low stresses soon after removal of a high stress. The dependence of

this mobility on the temperature could be an indicator of the DBTT.

3.3 Temperature Effects

Some of the physical properties investigated in this program could be measured as a function

of decreasing temperature through the region of the DBTT. These were Young's modulus, its associated damping capacity Q'g , and the magnetic parameters A,„ and ^t- Graphs of these

quantities as a function of temperature can be found in Section 2.3 for the elastic properties and

Section 2.7 for the micromagnetic properties. The magnetic hysteresis properties, H^, and B^ were measured at both liquid nitrogen temperature and at room temperature and showed no

unexpected variations. In general, all the properties varied smoothly between room temperature

and - 100 °C with no unusual behavior at the DBTT for the alloys involved. There was a sharp

peak in the internal friction near - 10 °C in the solution treated A710 steel which might be due to a dislocation relaxation phenomenon but relating it to the DBTT will require considerable

additional effort. The micromagnetic parameters were measured as a function of temperature on

the A533B steel samples cut from the weld and from the bulk of the SNUPPS RPV. At room

temperature, the measurements could not distinguish between the sources of the material.

However, at - 100° C the material from the weld (Sample Wl) and from the bulk (Sample 1-2)

could be distinguished by their values of >\,„ and jlI; . (See Fig. 2.7.12). Again, considerable more

study is needed to determine whether these observed differences in magnetic properties can be

related to the difference in DBTT for these two samples of A533B steel.

91
Table 3.1.2 Values for the mechanical strength parameters for the steels studied here.

Sample Condition Prop. Lim Upper YS 0.2% YS Ultimate Elong. Hardness 41JDBTT

MPa MPa MPa MPa % Rock. A °C

A710-CA Soln treat

A710-C5 Peak aged

A710-C7 Overaged

1508-AA Quenched

1508-A5 AsRec'd.

1508-A6 lh(a)675

173 — 415 650 28.8 55.5 -33

391 — 551 718 25.1 58.5 -10

446 471 471 558 29.2 53.2 -60

285 — 575 821 6.6 59 —

576 576 560 610 24.8 54 —

524 524 514 594 25.3 53 _

Table 2.1.3 Values of the mechanical strength parameters for the SNUPPS A533B.

Sample Condition Prop. Lim Upper YS 0.2% YS Ultimate Elong. Hardness 41JDBTT

MPa MPa MPa MPa % Rock. A °C

A533B 1-1 479 491 465 610 38 55 -20

A533B 1-2 461 468 452 574 43 53 -60

A533B 2-1 476 482 475 600 40 55 -25

A533B 2-2 513 525 507 640 35 56 -20

A533B W-1 568 589 578 672 24 57 -30

92
9i

•*^R

O

U

s

ena

9i

a

o

u

a,

s

'a

S

o

f<5

H

CO.

2 Qi.

= <

^ o

<0 X OS

<0 X

o

c

co

Q.E

V5

— ro

O o

o

o

— in

o —

o o

m — O O

O O

o <N O CM r~

r-j (N -^

in \c _^

:£ 2 t^

rr o —

(N r- t^

r~j — —

o o o

C3^ O O o o

inO

(N <N CNl (N (-s| (N (N (N

r--

t/1 OS «n o

VD '^ fN o (N fN

OO m

rOO o

r-i — —

r- •* r^* <n so

rsi

\0 sD so SO so

r~ 00 OS 00 o OS o o o o o Csl Csl Csl rsi fN <N o — —

^, t3 •a n ea u H, u

9 M OJO j3

ca ea o c

> c

1)3

00 o, U u-

< vn r- <f U U O < o o o QG

o

oU

rsi <

<

a:< ©

— CQ

sO<

^ "* I VI ># 1 >* 1

CQ<CQ

oo

93
C7

800

700

600

CA C5

500

400

300

200

100

X

^ ^ ^ HK' "

UTS

X— "

*" ..—""^

C3 , . cr "^ 'd'2%YS -^

^

"^ -v^ Prop. Lim.

53 54 55 56 Hardness (RA)57 58 59

Fig. 3.1.1 Relationship between three different measures of strength (the proportional limit, the

0.2% yield strength,and the ultimate tensile strength) and the hardness as measured on the

Rockwell A scale.

C7 CA C5

1.15

1.1

1.05

0.95

2 0.9

a.

0.85

0.8

0.75

Trans. Inc. Perm..

53 54 55 56 Hardness (RC A)

57 58 59

Fig. 3.1 .2 Physical properties that exhibit a monotonic change with hardness in the A710 alloy

series. For convenience in presentation, the properties have been normalized by their values in the

solution treated state (CA).

94
1

-r^ Proportional Limit C5 C7 ^"si^ ^""• — — — .^ ^__^ -A^ He

- - _ ^

0.9 . ^^k^

^X^^ Shear Mo^ ^ \

^V,,^ X ^ ^^>>^ X \

? °'® ^^^^ X \

^^^^^ ^\

jg ^V,^^ "^\

1 ^V^ g 0.7 ^V^^ ^

i

^s^,^^

Q. 0.6 ^^3- "^Log. Dec.

0.5 H@max N.

0.4 1 1 ^

150 200 250 300 350

Proportional Limit (MPa)

400 450

Fig. 3.1.3 Physical properties that exhibit a monotonic change with the proportional limit for the

A710 series. For convenience of plotting, the properties have been normalized by their values in

the solution treated state (CA).

S

1

•si

i

?

&

2

a.

C7

DBTT

CA C5

Int. Strain-

^~T

^^^0^""''^''^ 1.35

^_^,„*-'-'''''''^ Lambda

^^^^''^'^^^^'^'''''

* '^ 1.3 ^ ^ * frans. Inc. Perm. 1.25 / ^ -^ 1.2

1.15

W^ 1

i 1 1 1 1

1.1

1.05

1-

-60 -50 -40 -30 DBTT (deg. C)

-20 -10

Fig. 3.1.4 Relationship between four of the physical properties and the ductile-to-brittle transition

temperature (DBTT) for the three heat treatments of the Fe-1.13% Cu alloy. The properties have

been normalized by their values in the overaged condition (C7).

95
4. CONCLUSIONS

4.1 Task I

Task I was to define and investigate nondestructive physical property measurements that

could be promising candidates for predicting the DBTT in radiation embrittled A533B steel. The

most significant findings were:

1

.

Based on the measurements carried out in Task II, several physical properties appear

capable of predicting the ultimate tensile strength, the hardness, and the DBTT of a steel in which

the embrittlement mechanism is similar to that responsible for embrittlement of RPVs. These

properties are the nonlinear ultrasonic parameter p, the mximum transverse incremental

permeability \i^, and the high field average magnetostriction A°°. All three properties can be

deduced from ultrasonic measurements that can be performed nondestructively. Two other

properties, the internal friction Q"' and the field at which the magnetostriction is a maximum H^^^^ appear to be related more closely to the proportional limit than to the hardness or the DBTT.

2. Measurement of several physical properties including Young's modulus, the internal

friction parameter Q"'

, [i^ and A,°° were made as a function of temperature over a range that

included the DBTT in both the A710 and A533B alloys. All the properties measured under these

conditions showed a nearly linear temperature dependence with no special variations in the

vicinity of the DBTT.

3. The static stress needed to initiate plastic flow (that is, to produce dislocation breakaway

and motion) was near the proportional limit observed in conventional constant strain rate tensile

tests. Once deformed, however, the static stress needed to cause a release of dislocations and an

increase in internal friction was much lower than the proportional limit. Time-dependent

recovery of this internal friction at a low load was similar to that observed during recovery of

cold-worked material at elevated temperature.

4. The magnitude of the internal friction was so small (even after small amounts of plastic

deformation) that it would be difficult to measure by conventional ultrasonic pulse-echo

techniques.

4.2 Task II

Task II was to accurately measure physical properties on an A710 steel (Fe-1.13%Cu)

surrogate for the RPV steel in which tempering could be used to produce samples strengthened to different levels by the formation of copper rich precipitates (CRPs) similar to those suspected of

causing radiation induced embrittlement in A533B alloys. Three sets of samples were produced in

the solution treated, peak aged and overaged conditions. Tensile tests of constant strain rate and

Charpy V-notch tests were used to determine five measures of mechanical strength. Ultrasonic,

magnetic, and x-ray techniques accurately measured 10 physical properties. Correlations between

96
mechanical and physical properties fell into two classes: those that were related to large

deformation mechanical properties such as the ultimate tensile strength, and those that related to the initiation of plastic flow such as the proportional limit. The results can be summarized as follows:

1Three properties that exhibit a monotonically increasing correlation with the UTS, the

hardness and the DBTT are the high field average magnetostriction A,~, the maximum in the

transverse incremental permeability, ]l^ , and the nonlinear ultrasonic coefficient P that quantifies

harmonic generation. All three can be deduced from nondestructive ultrasonic measurements.

2. These three physical properties were found to be proportional to internal strains measured

by the broadening of x-ray diffraction peaks and presumed to arise from coherency strains

surrounding the CRPs.

3. Four other properties exhibit monotonically decreasing correlations with the proportional

limit and the 0.2% yield stress. These were the internal friction Q"', the field H^^x' at which the

magnetostriction is a maximum, the elastic moduli, and the coercive field, H^,. The latter two are much less sensitive to the proportional limit and 0.2% yield strength than are the former two.

4. Because the internal friction depends on irreversible depinning and motion of dislocations

after the proportional limit has been exceeded, this physical property may be unsuitable as a nondestructive test. However, the temperature dependence of the internal friction and its response to incremental changes in stress may be good indicators of the mechanical properties of

the material. Unfortunately, the magnitude of the internal friction for the A710 and A533B steels

was so small that it will be difficult to measure with conventional pulse echo techniques.

4.3 Task III

Tasks I and II produced a list of ultrasonic and magnetic properties that were sensitive to the microstructural features that control the mechanical strength of an A7 1 steel and an the

A533B steel. Task III was to discuss the possible ways in which these properties could be

measured through the stainless steel cladding that covers the inner surface of an actual pressure

vessel. Of the 10 measurables investigated, 3 appear able to meet the requirements of being

nondestructive, having good sensitivity to the desired mechanical property and being able to operate through the cladding. These are:

1Deducing quantities related to H^^^ and A,°° from measurements of the amplitude of a low

frequency shear wave generated by magnetostrictive coupling at the cladding-to-A533B steel interface.

2. Deducing a quantity related to P from measurements of the amplitude of a longitudinal

wave produced by mixing two shear waves.

97
3. The magnetic coercive force H^ of the A533B material under the cladding could be

deduced from measurements of the static magnetic field distribution around an electromagnet

placed on the cladding but this particular physical property (like the elastic moduli) does not

appear to be particularly sensitive to the CRPs responsible for embrittlement.

A more detailed discussion of making measurements through the cladding can be found in

Section 2.9 above.

5. RECOMMENDATIONS

1

.

Perform the physical property measurements that are related to the ductile-to-brittle

transition temperature of the A710 alloy surrogate or a second surrogate material whose

mechanical properties more nearly match the response of the A533B alloy to radiation damage.

2. Develop procedures for measuring those physical properties that are sensitive to the

DBTT on half compact tension specimens in a hot cell. Perform these measurements on

surveillance specimens that have been subjected to different neutron fluences in a reactor. Determine the quality of correlations between the measured physical properties and the DBTT

and the fracture toughness of the surveillance specimens.

3. Determine the accuracy and the reliability for measuring magnetostriction and nonlinear

acoustic parameters of the A533B alloy under the cladding.

98

U.S. GOVERNMENT PRINTING OFFICE 1998-0-673-0 18/40 104
NISTTechnical Publications

Periodical

Journal of Research of the National Institute of Standards and Technology—Reports NIST research

and development in those disciplines of the physical and engineering sciences in which the Institute is active. These include physics, chemistry, engineering, mathematics, and computer sciences. Papers cover a broad range of subjects, with major emphasis on measurement methodology and the basic

technology underlying standardization. Also included from time to time are survey articles on topics

closely related to the Institute's technical and scientific programs. Issued six times a year.

Nonperiodicals

Monographs—Major contributions to the technical literature on various subjects related to the Institute's scientific and technical activities. Handbooks—Recommended codes of engineering and industrial practice (including safety codes)

developed in cooperation with interested industries, professional organizations, and regulatory bodies.

Special Publications—Include proceedings of conferences sponsored by NIST, NIST annual reports, and

other special publications appropriate to this grouping such as wall charts, pocket cards, and

bibliographies.

Applied Mathematics Series—Mathematical tables, manuals, and studies of special interest to physicists,

engineers, chemists, biologists, mathematicians, computer programmers, and others engaged in scientific and technical work.

National Standard Reference Data Series—Provides quantitative data on the physical and chemical

properties of materials, compiled from the world's literature and critically evaluated. Developed under a worldwide program coordinated by NIST under the authority of the National Standard Data Act (Public

Law 90-396). NOTE: The Journal of Physical and Chemical Reference Data (JPCRD) is published bi- monthly for NIST by the American Chemical Society (ACS) and the American Institute of Physics (AIP).

Subscriptions, reprints, and supplements are available from ACS, 1155 Sixteenth St., NW, Washington,

DC 20056.

Building Science Series—Disseminates technical information developed at the Institute on building

materials, components, systems, and whole structures. The series presents research results, test methods, and performance criteria related to the structural and environmental functions and the durability

and safety characteristics of building elements and systems.

Technical Notes—Studies or reports which are complete in themselves but restrictive in their treatment of a subject. Analogous to monographs but not so comprehensive in scope or definitive in treatment of

the subject area. Often serve as a vehicle for final reports of work performed at NIST under the sponsorship of other government agencies.

Voluntary Product Standards—Developed under procedures published by the Department of Commerce

in Part 10, Title 15, of the Code of Federal Regulations. The standards establish nationally recognized

requirements for products, and provide all concerned interests with a basis for common understanding of

the characteristics of the products. NIST administers this program in support of the efforts of private- sector standardizing organizations.

Consumer Information Series—Practical information, based on NIST research and experience, covering

areas of interest to the consumer. Easily understandable language and illustrations provide useful background knowledge for shopping in today's technological marketplace.

Order the above NIST publications from: Superintendent of Documents, Government Printing Office,

Washington, DC 20402.

Order the following NISTpublications—FIPS and NISTIRs—from the National Technical Information

Service, Springfield, VA 22161.

Federal Information Processing Standards Publications (FIPS PUB)—Publications in this series collectively constitute the Federal Information Processing Standards Register. The Register serves as the

official source of information in the Federal Government regarding standards issued by NIST pursuant to the Federal Property and Administrative Services Act of 1949 as amended. Public Law 89-306 (79 Stat. 1127), and as implemented by Executive Order 11717 (38 FR 12315, dated May 11, 1973) and Part 6 of

Title 15 CFR (Code of Federal Regulations).

NIST Interagency Reports (NISTIR)—A special series of interim or final reports on work performed by

NIST for outside sponsors (both government and non-government). In general, initial distribution is handled by the sponsor; public distribution is by the National Technical Information Service, Springfield,

VA 22161, in paper copy or microfiche form.
U.S. Department of Commerce

National Institute of Standards and Technology

325 Broadway

Boulder, Colorado 80303-3328

Official Business

Penalty for Private Use, $300