NBSIR 83-2685

Effects of Variations in Heat

Treatment on the fVlechanica!

Properties and Microstructure of

ASTiV! A710 Grade A Class 3 Steel

U.S. DEPARTMENT OF COMMERCE

National Bureau of Standards

National Measurement Laboratory

Fracture and Deformation Division " Washington, DC 20234

May 1 983

Sponsored by

David Taylor Naval Ship Research & Development Center

Annapolis, MD 21402

NBSIR 83-2685

EFFECTS OF VARIATIONS IN HEAT

TREATMENT ON THE MECHANICAL

PROPERTIES AND MICROSTRUCTURE OF

ASTM A710 GRADE A CLASS 3 STEEL

G. E. Hicho

L. C. Smith

S. Singhal

R. J. Fields

U.S. DEPARTMENT OF COMMERCE

National Bureau of Standards

National Measurement Laboratory

Fracture and Deformation Division

Washington, DC 20234

May 1 983

Sponsored by

David Taylor Naval Ship Research & Development Center

Annapolis, MD 21402

U.S. DEPARTMENT OF COMMERCE, Malcolm Baldrige, Secretary

NATIONAL BUREAU OF STANDARDS, Ernest Ambler, Director
I

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I

i r

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Effects of Variations in Heat Treatment on the Mechanical Properties and

Microstructure of an ASTM A710 Grade A Class 3 Steel

G.E. Hicho*, L.C. Smith*, S. Singhal+, and R.J. Fields*

National Bureau of Standards

Washington DC 20234

A final report assessing the effects of variations in heat treatment on

the mechanical properties and microstructure of ASTM A710 grade A class 3

steel is given. The goal of this project was to evaluate the sensitivity of

the mechanical properties of A710, an age hardening steel, to heat treatment

variations that may occur in the production of thick plate. Tensile and

impact tests were conducted on the material heat treated under the various

conditions. Microstructure evaluations using optical and electron

metal lographic methods were conducted. Small angle neutron scattering

examinations were also performed on selected specimens in order to evaluate

the precipitate distribution. Results indicate that austenitization and aging

time has little effect on the resultant mechanical properties. However, the

mechanical properties were found to be sensitive to aging temperature. Impact

results at 0°F indicate that aging at 900°F is detrimental in that the

fracture surfaces of the broken impacts exhibited cleavage. Furthermore, the

results suggest that welding or similar heating operations will not seriously

degrade this material's properties.

Key Words: age hardening steel, cleavage fracture, ductile fracture, energy

absorbed, lower yield strength, microstructure, precipitate, ultimate tensile

strength, upper yield strength.

*Fracture and Deformation Division, Center for Materials Science, National

Measurement Laboratory

+Reactor Division, Center for Materials Science, National Measurement

Laboratory

1
ADMINISTRATIVE INFORMATION

The investigation reported herein was funded by the Submarine Materials

Block Program sponsored by Dr. H. H. Vanderveldt, Naval Sea Systems Command

(SEA 05R25), and funded under Program Element 62761N, Task Area SF-61-541-591

,

Work Unit 2803-132. The effort was monitored by Mr. E. J. Czyryca, David

Taylor Naval Ship R&D Center (Code 2814). This report satisfied in part^

Milestone PS 3.1/2 of Submarine Block Task, "Development of Thick-Plate HSLA

Steels for Naval Applications", which is supervised by Mr. J. P. Gudas, David

Taylor Naval Ship R&D Center (Code 2814).

2
INTRODUCTION

This report presents data as to whether temperature and time variations

in the recommended heat treatment have any effect on the mechcnical properties

and microstructure of ASTM A710 grade A, class 3 steel. The recommended heat

treatment requires the steel to be austenitized at 1650°F for 68 minutes,

water quenched, and aged at 1100°F for 30 minutes followed by air cooling.

In this report, results will be shown which compare the mechanical

properties obtained after short (30 minute) and long (90 minute)

austenitzation times at 1650°F, followed by aging times of 30,60, and 90

minutes at temperatures of 900, 1000, 1100, and 1200°F.

The steel, is an age-hardening ferritic steel containing about 1.2 weight

percent copper. The copper is the primary strengthening element. Basically,

the strengthening mechanism is as follows. During a soaking period at the

austenitization temperature, a certain amount of copper goes into solid

solution. When the steel is water quenched, the copper is retained in a

supersaturated solid-solution. Subsequent aging of the steel at a selected

temperature and time allows fine copper-rich precipitates to form within the

ferrite. These precipitates impede dislocation motion, hence controlling the

yield strength of the steel. In the as-quenched condition, the yield strength

is about 60KSI. By aging this steel at 1100°F for 30 minutes, the recommended

treatment, the yield strength is increased to 80-90KSI. In addition to the

strengthening brought about by the precipitation of the fine copper phase,

further strengthening is obtained by the addition of Cb. Columbium, when

added to the melt, enhances grain refinement in the steel. This addition

tends to lower the y-a transformation, thereby enhancing nucleation, and at

the same time retarding the growth of ferrite. This leads to a very fine

ferrite grain size. Because precipitation hardening is the primary

3
strengthening mechanism, the melt carbon content can be low, thereby promoting

good weldability, while retaining good ductility and strength. Since copper

is one of the primary alloying elements in this steel, it is quite possible

that hot shortness could be a problem. To eliminate this problem, nickel is

added to the steel (1). Small amounts of Cr and Mo are also added to retard

auto-aging which might occur after hot rolling. The combination of low

carbon, strengthening by aging of the epsilon copper phase, and grain

refinement appears to produce a steel with significantly improved strength and

toughness qualities. (2) Jesseman et. al. and Hurley et. al . (3) have

published papers showing the mechanical properties of the steel after

austenitizing and aging at the recommended temperatures.

Optical examinations were conducted on the test coupons in order to

obtain a correlation between microstructure and strength. Scanning electron

microscopy (SEM) examinations were also conducted on the fracture surfaces of

both the tensile and impact specimens. Small angle neutron scattering (SANS)

studies were conducted on some samples as a function of aging time and

temperature in order to evaluate both copper precipitate size and

distribution. Correlations were found between the ultimate tensile, upper

yield, lower yield, impact energy, and hardness. With these results,

conclusions are drawn as to the sensitivity of A710 to variations in heat

treatment.

AS RECEIVED PLATE

A plate of A710 steel, approximately 1 1/4 inch x 2 feet x 2 feet was

used in the investigation. The mill report indicates that the plate was

austenitized at 1650°F for 68 minutes and then water quenched. The

4
plate was subsequently aged hardened at 1100°F for 30 minutes and air cooled.

The rolling direction was marked on the plate with paint.

Before test coupons were removed from the plate, a sample was removed and

metal lographi cal ly examined to verify the marked rolling direction. It was

hoped that the inclusion shape and orientation would reveal the rolling

direction. Figure 1 shows the unetched steel in the three orthogonal

directions. The photomicrographs reveal the presence of globular sulfides

indicating that inclusion shape control was used (usually by Ca addition to

the melt). When low sulfur contents (.010 weight percent or less) are

requested, Ca is added to the steel. Ca additions tend to globularize the

sulfur in the steel and this dramatically improves the shelf energy of the

steel. Because of the Ca addition, it was almost impossible to ascertain the

plate's rolling direction, hence the direction marked on the plate was taken

as the rolling direction. The chemical composition, as reported by the

-manufacturer is given in Table 1. Of note is the very low sulfur content

(.004 weight percent).

MICROSTRUCTURE OF THE AS RECEIVED PLATE

The etched microstructures , in an orthogonal representation, are shown in

Figure 2. The microstructure consists mainly of massive ferrite which,

despite its name, is very fine grained. In the interior of the plate, there is

evidence of some long range diffusional decomposition product. This is due

probably to the fact that the interior of the plate can not be cooled as

rapidly as its surface. Optical photomicrographs of the plate at higher

magnification. Figures 3a, b, and 4a, b, show the presence of pearl ite and

possibly some bainite. The dull appearance of the ferrite could be

due to the presence of an excessive amount of copper (4). It is

5
quite possible that some retained austenite is present, but this possibility

was not investigated further.

The mill report stated that the grain size number was 8; however, our

measurements found the grain size to be number 9.5 to 10. This finer grain

size is not detrimental. Quite the contrary, the finer grain size enhances

the mechanical properties of the steel.

A high resolution electron micrograph of A710 is shown in Figure 5. The

dark particles are the incoherent FCC copper-rich precipitates. No finer

precipitates were found. As noted by Hornbogen (5), there is no source of

diffraction contrast for coherent copper precipitates in iron. For this

reason, neutron scattering techniques were employed and are detailed in

Appendix A.

SECTIONING OF THE TEST COUPONS FROM THE AS RECEIVED PLATE

The as-received plate was cut along the lines shown in Figure 6. Each

test coupon was approximately 6 in. long, 2 1/2 in. wide and the thickness of

the plate. From each coupon, two .500 in. diameter tensile specimens and

three standard Charpy V-notch (CVN) specimens were prepared. The specimens

were taken as close as possible to the quarter-thickness location in the

plate. The tensile specimens were oriented with the longitudinal axis

transverse to the rolling direction. The orientation of the CVN specimens

corresponded to the ASTM TL direction; that is, the axis was transverse to the

rolling direction of the plate and the notch located parallel to the rolling

direction. It should be noted that each individual heat treatment was

performed on each coupon prior to machining.

HEAT TREATING OF THE TEST COUPONS

The heat treatment schedule is shown in Table 2. Before heat treating

each coupon, the time to reach the austenitization temperature of 1650°F, and

6
the times to reach the aging temperatures of 900°F, 1000°F, 1100°F, and 1200°F

were determined. To accomplish this, thermocouples were inserted into

selected coupons and temperature/time curves. Figure 7, for each condition

were obtained. In addition, a thermocouple was also immersed into the quench

medium in order to monitor the bath temperature which was maintained at 70°F.

PREPARATION OF TEST SPECIMENS

The specimens used for the tensile tests were .500 inch in diameter and

dimensioned according to ASTM Specification A370-77,! (unchanged through

1982). The Charpy V-notched specimens were prepared and tested according to

ASTM Specification E23-82. The notches were placed in the specimen with a

broaching tool for reproducabil ity.

TEST CONDITIONS

The tensile specimens were tested at room temperature using a universal

testing machine. Calibration data for the testing machine indicated that the

accuracy of the load readings was within ±1% as recommended by ASTM Committee

E-4 on testing. An extensometer was placed on the specimen to measure the

percent strain over a 2 inch gage length. Figure 8 shows a typical

stress/strain curve obtained for this material. ASTM specification A370-82

recommended the loading rate be O.Olin/in/min prior to yield, then

O.lin/in/min after yield for a specimen having a gage length of 2 inches.

The CVN specimens were tested at 0°F according to ASTM specification

, E23-82. They were immersed in the bath medium (ethyl alcohol) and held at

temperature for at least 30 minutes before testing. The bath was magnetically

stirred and the temperature was constantly monitored with a thermocouple.

7
TEST RESULTS (Summary Tables)

A summary of the test results (tensile, impact, and hardness) for the

specimens austenitized at 1650°F for 30 minutes, water quenched, and aged at |

the appropiate temperatures and times are shown in Table 3. Similarly,

results for specimens austenitized at 1650°F for 90 minutes, water quenched,

and aged are shown in Table 4. In addition, tests were conducted on the

as-received and the as-quenched plates. Specimens were also prepared after

heat treating at the recommended austenitization temperature time of 1650°F

for 68 minutes, and aged at 1100°F for 30, 60, and 90 minutes. These results

are shown in Table 5.

TENSILE TEST RESULTS

In Figure 9, both the ultimate tensile stress (UTS) and lower yield

stress (LYS)* increase when the material is aged for 30 minutes at 1000°F as

compared to 900°F. This phenomenon was not found at aging times greater than

30 minutes. At temperatures greater than 1000°F, irregardless of aging time,

both the UTS and LYS decrease drastically. Elongation and reduction of area

increase at temperatures above 1000°F because of internal metallurgical

mechanisms. The absence of data for specimen aged at 1200°F was this specimen

was inadvertantly aged at 1100°F.

The shape of the tensile data curves. Figure 10, for specimens

austenitized at 1650°F for 90 minutes, water quenched, and then aged, appear

to be quite similar to these shown in Figure 9. Again there is an increase in

the UTS and LYS as the aging temperature is increased from 900°F to 1000°F

and, like the specimens austenitized for 30 minutes the UTS and LYS both

decrease as the aging temperature increases. Elongation and the reduction of

II

*The LYS is equivalent to the 0.2% offset yield stress. I

8
area also follow the trend established before - both increase as the aging

temperature increases.

HARDNESS TEST RESULTS

Hardness results are shown in Tables 3, 4, and 5. Six Rockwell A

measurements readings were taken on each impact specimen before testing.

These values were then converted to the approximately equivalent Rockwell B

values for presentation in the tables.

Figure 11 shows hardness plotted as a function of precipitation hardening

time. Hardness decreases as both the aging time and temperature increase.

One exception to that trend occurs at the aging temperature of 900°F,

regardless of austenitizion time. At that aging temperature, the hardness

increases significantly after aging times of 60 and 90 minutes. This

indicates that the A710 is underaged at 900°F and overaged at all higher

temperatures including the temperature of the manufacturer's recommended

treatment. Further heating of overaged material, such as might occur in the

heat affected zone (HAZ), will only slightly decrease the hardness.

These results are more clearly shown in Figures 12 and 13 where hardness,

HRB is shown as a function of precipitation hardening temperature. It is

shown that at aging temperatures greater than 900°F, the hardness obtained

with longer aging times and higher temperatures is always lower that obtained

with shorter aging times and lower temperatures. At 900°F, however, the

hardness obtained after aging times of 60 and 90 minute was greater than that

obtained after 30 minutes. This appears to be due to the thermodynamics

associated with the epsilon-copper phase formation. Previous work on Fe-0.9%

Cu (5) has shown that after short aging times of 2, 8, or 10 minutes, at

900°F, hardness is still increasing, whereas for the same aging times, at

1000®F, the hardness has leveled off. It appears from the hardness data for

9
the A710 that at the aging temperature of 900®F, the maximum hardness has not

yet been achieved, even for the aging time of 90 minutes.

The curves in Figures 12 and 13 are quite similar to the strength versus

precipitation hardening temperature curves shown in Figures 10 and 11. Like

the strength curves, the highest hardness was obtained at the aging

temperature of 900°F and aging time of 90 minutes.

IMPACT TEST RESULTS

Charpy V-notch impact test results for specimens tested at 0°F are shown

in Tables 3 and 4 and Figures 14 and 15. The energy absorbed values are the

lowest after aging at 900°F, regardless of austenitization time. As the aging

temperature increases, so does the energy absorbed. An anomaly is observed

for the specimen austenitized for 90 minutes and then aged for 30 minutes at

1000°F. The energy absorbed is essentially the same as that obtained at an

aging time of 30 min at 900°F. The reason for this is unknown at this

writing. Comparison of Figures 14 and 15 indicates that there is more scatter

in the data for the material austenitized for 90 minutes.

Table 6 shows the results of the energy absorbed values obtained for both

austenitizing times and at the various aging temperatures. The 900°F

treatment, which is shown to result in low impact properties, corresponded to

the highest UTS and LYS values. It can also be seen that overaging results in

improved impact resistance. This has good implication with regard to weld

heat affected zone toughness. That is, further heating will only toughen the

material. From these impact results, it can be concluded that the

austenitization time has no significant effect on impact properties.

FRACTOGRAPHIC EXAMINATION RESULTS

A. Tensile Specimens

Photographs, taken at low and high magnifications of selected broken

tensile specimens, are shown in Figures 16 and 17.

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Figure 16 shows the normal cup and cone type fracture. Corresponding

scanning electron micrographs, Figure 17, taken at a higher magnification,

show fracture surfaces characteristic of a ductile material - voids that have

coalesced and subsequently fractured.

Upon examination of other fractured tensile specimens, splitting in the

fracture surface was observed. Some of the specimens that showed splitting

are shown in Figure 18. Splits in the material were found parallel to the

loading direction. A higher magnification observation using the SEM, Figure

19, taken along the split, was unable to clearly identify the fracture

characteristics of the split.

' In another SEM examination of a specimen that was severely split. Figure

20a, b, and c, it was observed that the fracture was predominately ductile in

nature as evidenced by the shear observed on the fracture surface. No further

conclusions could be drawn as to the cause of the splitting; however, it was

suggested that possible causes could have been stress induced hardening, an

orientation effect, or possibly temper embrittlement. We feel, however, that

splitting could be a stress state problem since two identically heat treated

specimens responded differently when tested. While the tensile specimen that

was machined to a diameter of 0.500 inch showed splitting, one that v;as

machined to 0.252 inch showed no splitting at all. This specimen showed cup

and cone fracture. The specimen that split did not show any evidence of

cleavage fracture when observed on the SEM.

Figures 21 and 22 are graphical representations of the type of fracture

observed on the test specimens as a function of the austenitization time and

precipitation hardening temperature.

The figures show that some degree of splitting was observed at aging

temperatures of 900°F, 1000°F, and 1100°F. At the aging temperature of

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1200°F, irregardless of aging time, there was no evidence of splitting

observed in the tensile specimens. Also at an austenization time of 90

minutes followed by an aging temperature and time of 900°F, 30 minutes

respectively, was there an absence of splitting. The length of the

austenitization time does not appear to dictate whether the specimen splits on

loading.

Even though this splitting phenomenon was observed in some tensile

specimens, it should be realized that the splitting occurred at reductions of

area near 76% and elongations of 26%, well into the plastic instability region

of the steel

.

B. Impact Specimens

The fracture surfaces of several impact specimens were examined using a

light microscope. There was no evidence of splitting in these specimens. A

low magnification photograph of several broken impact specimens is shown in

Figure 23. SEM fractographs of the same surfaces are shown in Figure 24. The

fracture appearance of these are definitely ductile in nature.

Impact specimens fractures corresponding to the tensile specimens that

contained splits were examined in the SEM. Some of these specimens showed

evidence of cleavage fracture, others did not. With these conflicting

results, it was decided to examine all of the impact specimen fractures with

the SEM in order to determine which ones exhibited cleavage fracture and which

ones did not. A micrograph of a impact specimen that revealed cleavage

fracture is shown in Figure 25a, b, and c.

Figures 26 and 27 show the energy absorbed versus precipitation hardening

temperature for both austenitization times. Printed on these curves are the

letters C and D. The letters C and D appearing in these figures correspond to

the observed fracture appearance - C indicates cleavage fracture and D

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ductile fracture. It was found that cleavage fracture was observed only on

those impact specimens that were aged at 900°F. Again, austenitization time

does not appear to affect whether the fracture was ductile or brittle in

nature.

For those specimens that showed cleavage, the cleavage did not extend

over the entire surface. It was located primarily in the vicinity of the

hammer's impact. The remainder of the sample's fracture surface was ductile

exhibiting dimpled rupture.

CORRELATIONS BETWEEN OBSERVED RESULTS

It was perceived that a correlation might exist among the mechanical

property results obtained in this work. One such correlation was developed

using a combination of the UTS, the LYS, and energy absorbed. Figure 28 is a

plot of these data.

In this figure, it is obvious that the correlation is not linear.

However, it is shown that with a known UTS or LYS, one can roughly estimate a

corresponding energy absorbed. The curves in Figure 28 also reveal that

impact properties improve as the UTS and LYS decrease. There is a rapid

decrease in these tensile properties once the impact energy exceeds 170

ft-lb.

Figure 29 is a plot of hardness and corresponding UTS and LYS values.

Again, a linear correlation was not found, but as in the previous correlation

for UTS, LYS, and energy absorbed, one could for only those heat treatments

studied, knowing the hardness determine with reasonable accuracy either the

UTS or LYS.

Figure 30 is a plot of hardness versus energy absorbed. From this curve

it is possible to determine the approximate impact value from the hardness

data. These data may be used to determine the most impact sensitive region of

a HAZ. 13
One noticeable tensile property characteristic of this material is that

it exhibits a load drop on yielding. This was shown in Figure 8, a drawing of

the typical stress/strain curve that was obtained. Of importance is the

degree of load drop; that is, the difference between the upper and lower yield

stress values. This is probably due to the pinning of mobile dislocations by

the copper precipitates. An analysis of the load drop for the specimens

tested follows.

Figures 31 and 32 are plots of the upper yield stress minus the lower

yield stress vs aging temperature, for different austenitization times. It is

shown in these figures that this difference is most pronounced for material

that was aged at 1100°F for austenitizing at either 30 or 90 minutes.

Figures 33 and 34 show curves developed by plotting the ratios of UYS and

LYS to UTS versus precipitation hardening tmeperature. The LYS/UTS ratio is

essentially constant over the range of aging temperatures, whereas the UYS/UTS

ratio is very sensitive to the aging temperature. This sensitivity most

pronounced is again at the reported optimum aging temperature, 1100°F. At

this temperature the UYS/UTS is almost 99% of the UTS. While this is not

significant in displacement controlled situations, it may be relevant in load

control. It should be noted that at the optimum aging temperature of 1100°F,

the reduction of area and elongation were found to be 75% and 26%

respecti vely.

SMALL ANGLE NEUTRON SCATTERING (SANS) RESULTS.

The SANS work is detailed in Appendix A. Results indicate that, at an

austenitization time of 90 minutes and aging temperature of 900°F for 30

minutes, the volume fraction of the epsilon-copper precipitate was lower

than that observed for the other heat treating conditions. However,

the distribution of the precipitate was finer, with closer inter

particle distances, which led to a high hardness. It was also shown

that the volume fraction of the precipitate increases as the aging
temperature increases. Furthermore, the excellent correlation between

hardness and the total neutron cross section for the isochronical ly aged

samples demonstrates the power of this microstructural tool.

DISCUSSION

In this work, the sensitivity of the mechanical properties A710 Grade A,

Class 3 steel, to variations in heat treatment were determined. In

particular, tensile and impact properties were measured as a function of

austenitzation time, and aging temperature and time.

A710 steel is a precipitation hardening steel, and the element which is

primarily responsible for this hardening is copper. After the proper

austenitization time and temperature, the yield strength can be raised from

its solution heat treated value of 60KSI to a yield strength in the range of

80-90KSI. This steel, used in the optimum quenched and aged condition,

exhibits a microstructure consisting mainly of massive ferrite, which despite

its name is very fine grained. The small grain size of the steel is

attributed to the Cb addition which forms a columbium-carbon-ni tride Cb(C,N)

that tends to suppress grain growth during austenitization. It also reduces

the transformation temperature. The measured grain size was 9.5 to 10 for the

plate tested.

An examination of the as-polished surface, in three orthogonal

directions, revealed the presence of globular inclusions. This is an

indication that Ca was added to the steel for the purpose of inclusion shape

control. Inclusion shape control (Ca addition) improves the impact properties

significantly, in particular the shelf energy.

Transmission electron microscopy studies of the received plate revealed

o large, > 200 A incoherent, copper precipitates that are probably responsible

15
for the increase in the yield strength on aging. The coherent epsilon-copper

precipitate was not resolved in the STEM because of its small scattering

contrast for electrons and because of its lack of coherency strain.

A more detailed analysis at the precipitation behavior was carried out by

a technique called small angle neutron scattering (SANS). It was demonstrated

(appendix A) that SANS could be used to determine the distribution and size of

the precipitates.

Following these examinations, as received 1 1/4" thick plate was

sectioned into coupons and each coupon heat treated according to a

pre-determined schedule. Tensile and impact specimens were sectioned from the

1/4 thickness location of the plate and subsequently tested.

Previous tensile test results clearly indicate that A710 steel is

significantly affected by the aging time but not the austenitizing time. The

tensile tests also revealed that A710 exhibits upper and lower yield points.

The change from the UYS to the LYS is called yield drop. This yield drop is

followed by Luders band propagation, which is in turn followed by work

hardening until the ultimate strength is reached. The ductility as measured

by the reduction of area and elongation in the tensile tests was about 75% and

26% respectively - both indicative of a very tough material.

Table 7 shows a comparison of the UTS, UYS, and LYS values shown in

Figures 9 and 10. It was reported by the manufacturer that the optimum aging

temperature was 1100°F. Test data for specimens aged at temperatures below

1000°F were not found. However, in our investigation, specimens were aged at

900°F and tested. It was found that the tensile properties obtained after

aging times of 60 and 90 minutes were greater than those reported for the

optimum aging temperature and time. These results appear to be due to the

size of the epsilon-copper precipitate formed after the 90 minute aging time

16
at 900F. The size of the precipitate formed is so fine and dispersed

throughout the ferrite that dislocation motion is severely retarded. This

dislocation retardation leads to higher flow stresses. At 1200°F, the

precipitates are no longer fine, but are coarse due to overaging. The

resulting precipitates are no longer coherent with the matrix and dislocations

motion is much less difficult. This is manifested in lower flow stresses.

The tensile results for the steel aged at various temperatures indicated

that the maximum strength was obtained after age hardening at a temperature of

900F, the impact energy values for this same aging temperature, tested at 0°F

were low. SEM observations of these fracture surfaces show some cleavage. At

an aging temperature of 1000°F and above, it appears that the material has

passed through an optimum and is in an overaged condition. This overaging is

advisable because overaging improves the impact properties significantly.

The 900°F aging treatment corresponds to the highest ultimate and yield

strength value, but lowest impact toughness. Overaging results in improved

impact resistance and has good implications in regard to the heat affected

zone weld toughness.

Attempts were made to develop correlations among tensile test values,

hardness, and impact results. The most significant relationships obtained

were those relating the tensile test results (UTS, LYS, UYS) and hardness.

Attempts were made to relate the observed metallurgical microstructure of the

aged material with the obtained mechanical properties. However, as shown in

Figure 35, the microstructure did not vary noticeable for different aging

conditions, hence this approach was disregarded. Because of this

17
observation, SANS was chosen. SANS (appendix A) did prove to be a viable tool

in quantifying the precipitate distribution as a function of aging time and

temperature. However the precipitate was not defined.

Scanning electron microscopic work on the fractured test samples revealed

that some of the tensile specimens showed splits on the fracture surface. The

splitting was thought to be due to the stress state. This splitting could

be accentuated if the tensile specimens were tested in the through thickness

direction. No through thickness specimens were tested in this work. The

impact specimens did not show any evidence of this splitting. A710 grade A,

class 3 steel was found to be insensitive to austenitization time, but quite

sensitive to aging time and temperature.

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CONCLUSIONS

1. Variations of ±30 minutes in austenitizing time at 1650 °F have very

little effect on mechanical properties.

2. Variations of ±100 ®F and ±30 minutes in the manufacturer' s recommended

aging treatment do not seriously degrade mechanical properties.

3. Tensile ductility is virtually unaffected by any of the variations noted

in 1 and 2 above.

4. Degraded mechanical properties can be restored by suitable

reaustenitization and aging treatment.

5. Manufactured plate is normally in an overaged condition. Further heating

of this material by welding or other operation should only increase impact

resistance.

6. Underaging at 900 °F considerably reduces impact properties.

7. UYS can approach UTS for some conditions. This may have ramifications

for load controlled situations.

8. Longitudinal splitting of tensile specimens and cleavage during impact can

occur for certain conditions. This may have ramifications for thicker

sections or welded structures.

19
REFERENCES

1. Hydrean, P.P., Kitchin A.L., and Schaller, F.W. , "Hot Rolling and Heat

Treatment of Ni-Cu-Cb(Nb) Steel," Met. Trans., Sept. 1971, 2, p. 2541.

2. Jesseman, R.J., and Murphy, G.J. "High Strength Precipitation Hardened

Alloy Steel For Heavy Section Application," Presented at the 1977 Triple

Engineering Shows and Conferences, Chicago, see. Oct. 26, 1977.

3. Hurley, J.L., and Shelton, C.H., "Age-Hardenable Nickel -Copper Steels",

Metals Quarterly, May, 1956, p. 25.

4. Creswick, W.E., "Commercial Development of A Rimmed Low-Alloy

Precipitation-Hardening High Strength Steel," British Iron and Steel

Institute, Publication 104, 1967.

5. Hornbogen, E., "Aging and Plastic Deformation of an Fe-0.9% Cu Alloy",

Trans. ASM, Vol . 57, 1964, p. 120.

20
Table 1. Chemical analyses* of as received plate.

Heat

Number 55180

Grain Size 8

Element

Weight

Percent

C 0.04

Mn 0.58

P 0.010

S 0.004

Si 0.30

Cr 0.68

Ni 0.87

Mo 0.19

Cu 1.20

Cb 0.046

* As reported by manufacturer.
Table 2. Heat treating schedules for test specimens.

Short Austenitizing Time Long Austenitizing T ime

1650 °F - 30min - W/Q 1650°F - 90min - W/Q

Specimen Aging Aging Specimen Aging Aging

Number Time (mi n)Temp °F Number Time(min)Temp

1 30 900 13 30 900

2 30 1000 14 30 1000

3 30 1100 15 30 1100

4 30 1200 16 30 1200

5 60 900 17 60 900

6 60 1000 18 60 1000

7 60 1100 19 60 1100

8 60 1200 20 60 1200

9 90 900 21 90 900

10 90 1000 22 90 1000

11 90 1100 23 90 1100

12 90 1200 24 90 1200

A1 so :

1650F - 68 min - W/Q then

Specimen 25 30 min @ 1100°F

" 26 60 min @ 1100°F

" 27 90 min @ 1100°F

28 As quenched, no aging

29 As received.

22
Table 3. Mechanical property results for specimens austenitized at 1650°F for

30 minutes, water quenched and aged at the indicated temperatures

and times. Impact specimens tested at 0°F.

Specimen Aging Aging UTS^ UYS^

Number Time (min) Temp°F (KSI) (KSI)

1 30 900 109.2 99.1

2 30 1000 111.5 104.8

3 30 1100 97.8 95.6

4 30 1200 90.3 89.4

5 60 900 113.2 104.5

6 60 1000 109.2 101.4

7 60 1100 96.3 95.8

8 60 1200 89.3 87.3

9 90 900 114.3 102.4

10 90 1000 105.7 97.2

11 90 1100 94.5 92.5

12 90 1200 NOT

LYs^»* EL®

(KSI)(%in2")

RA^

(*)

hard‘s

(HRB)

ENERGY^

ABSORBED

(Ft-lb)

96.1 25.9 73.9 98.5 128.7

98.0 26.6 72.3 99.0 161.3

86.3 27.3 76.8 95.3 178.3

81.5 29.4 79.5 92.0 195.0

101.3 26.0 73.0 99.8 107.3

95.7 25.9 72.8 98.6 155.0

86.2 28.3 76.4 95.0 178.3

80.3 29.5 79.1 91.6 189.3

100.4 25.9 72.5 HRC 23.4 88.3

91.4 27.3 73.3 98.0 156.7

84.3

TESTED

28.2 77.8 94.4 185.7

(a) Average of 2 values

(b) Average of 6 values

(c) Average of 3 values

* LYS is equivalent to 0.2% effect yield stress.

23
Table 4. Mechanical property results for specimens austenitized at 1650°F for

90 minutes, water quenched and aged at the indicated temperatures

and times. Impact specimens tested at 0°F.

. ENERGY^

Specimen

Number

Aging Aging UTS

Time(min) Temp°F (KSI)

UYS°

(KSI)

LYS°

(KSI)

EL°

(%in2"

RA°

) {%)

HARD^

(HRB)

ABSORBED

(Ft-lb

13 30 900 109.2 101.0 96.4 24.9 74.5 98.7 129.7

14 30 1000 111.7 103.4 98.1 27.3 72.7 99.8 128.7

15 30 1100 97.9 95.4 87.4 27.8 76.8 95.6 178.0

16 30 1200 90.4 88.7 81.0 29.7 79.4 91.2 193.0

17 60 900 110.9 101.0 97.7 25.4 72.0 99.8 96.3

18 60 1000 109.0 103.7 96.8 25.2 71.4 98.3 168.3

19 60 1100 95.4 93.2 85.4 27.8 76.8 94.4 178.0

20 60 1200 88.2 87.2 78.9 28.9 78.5 90.4 213.7

21 90 900 115.4 107.9 103.7 25.8 71.8 HRC 23. 2 89.3

22 90 1000 105.8 100.1 93.8 26.0 73.4 98.5 167.0

23 90 1100 94.7 93.6 84.9 28.8 74.5 93.8 183.7

24 90 1200 85.2 83.5 76.6 30.8 80.8 89.0 204.0

(a) Average of 2 values

(b) Average of 6 values

(c) Average of 3 values

* LYS is equivalent to 0.2% effect yield stress.

24
Table 5. Results for specimens austenitized at 1650°F for 68 mi

water quenched. Impact specimens tested at 0°F.

Specimen Aging Aging UTS^

Number Time(min) Temp°F (KSI)

25 30 1100 99.4

26 60 1100 95.9

27 90 1100 92.3

28 As Quenched

(No Aging)

98.9

29 As Received 95.2

Mill Report Test

(1)

97.6

(2)

93.7

(a) Averagy of 2 Values

(b) Average of 6 Values

(c) Average of 3 Values

(d) 0.2% offset

* LYS is equivalent to 0.2%

UYS®

(KSI)

a LYS

(KSI)

EL^

(%in2"

RA^

){%)

hard‘s

(HRB)

98.3 89.3 28.3 75.3 95.5

94.4 86.1 27.7 76.9 94.6

89.8 82.8 28.7 77.8 92.6

57.5^ 27.8 72.9 94.2

90.8 81.7 28.3 76.1 93.8

89.0^ 25.0

87.2° 26.0

effect yield stress.

and then

ENERGY^

ABSORBED

(Ft-Lb)

184.0

190.7

189.7

169.3

200.0

25
Table 6. Comparison of Charpy V-notch impact test results for both

austenitization times. Specimens tested at 0®F.

Aging

Time(min)

Aging

Temp°F

1650°F - 30 min-W/Q 1650°F - 90 min-W/Q

Energy Absorbed*

(Ft - Lb)

Energy Absorbed*

(Ft - Lb)

30 900 128.7 129.7

30 1000 161.3 128.7

30 1100 178.3 178.0

30 1200 195.0 193.0

60 900 107.3 96.3

60 1000 155.0 168.3

60 1100 178.3 178.0

60 1200 189.3 213.7

90 900 88.3 89.3

90 1000 156.7 167.0

90 1100 185.7 183./

90 1200 (Not Tested) 204.0

Also 1650°F - 68 min. - W/Q

30 1100 184.0

60 1100 190.7

90 1100 189.7

As Quenched (No Aging) 169.3

As Received 200.0

*Average of 3 Values

26
Table 7. Comparison of UTS, UYS, and LYS results.

Aging

Time(min)

Aging

Temp°F

1650°F - 30 min - W/Q 1650°F - 90 min - W/Q

UTS

(Ksn

UYS

(KSI)

*' LYS

(KSI)

UTS

(KSI)

UYS

(KSI)

* LYS

(KSI)

30 900 109.2 99.9 96.4 109.2 101.0 96.4

30 1000 111.5 104.8 98.0 111.7 103.4 98.1

30 1100 97.8 95.6 86.3 97.9 95.4 87.4

30 1200 90.3 89.4 81.5 90.4 88.7 81.0

60 900 113.2 104.5 101.3 110.9 101.0 97.7

60 1000 109.2 101.4 95.7 109.0 103.7 96.8

60 1100 96.3 95.8 86.2 95.4 93.2 85.4

60 1200 89.3 87.3 80.3 88.2 87.2 78.9

90 900 114.3 102.4 100.4 115.4 107.9 103.7

90 1000 105.7 97.2 91.4 105.8 100.1 93.8

90 1100 94.5 92.5 84.3 94.7 93.6 84.9

90 1200 (Not Tested) 85.2 83.5 76.6

* LYS equivalent to 0.2% offset yield stress.

27
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STRESS

STRAIN

Figure 8. Typical stress/strain curve obtained for a tensile-test specimen.

35

Figure 9. Tensile test results for specimens austenitized at 1650°F for

30 minutes and then aged at the indicated temperature.

36

500“C 550°C eoo^c 650°C

O)

cn

CO

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116

no —

102 -

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o

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Figure 10. Tensile test results for specimens austenitized at 1650°F for

90 minutes and then aged at the indicated temperature.

37
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HARDNESS,

HRB

30 60 90

PRECIPITATION HARDENING TIME , MIN

Figure 11. Hardness, HRB, versus precipitation hardening time in minutes.

38
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500®C 550«C 600*C 650®C

PRECIPITATION HARDENING TEMPERATURE

Figure 12. Hardness, HRB, versus precipitation hardening temperature for

A710 austenitized at 1650°F for 30 minutes and then aged at the

indicated temperature and time.

39
0*05 & D®oo5 6;

-- t-..-i» | « I III I I win - I

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500®C 550®C 600°C 650®C

PRECIPITATION HARDENING TEMPERATURE

Figure 13. Hardness, HRB, versus precipitation hardening temperature for

A710 austenitized at 1650°F for 90 minutes and then aged at the

indicated tmeperature and time.

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500«C 550®C 600*C 650*C

270

250

230

210

190

170

150

130

110

PRECIPITATION HARDENING TEMPERATURE

Figure 14. Energy absorbed versus precipitation hardening

temperature for A710 austenitized at 1650°F

for 30 minutes and then aged at the indicated

temperature and time.

41

JOULES
'-aS Tli

Cifi

on

0NW-ii»w£>r-3?

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ENERGY

ABSORBED,

FT

-LB

500®C 550®C 600«C 650«C

320

310

290

270

250

230

210

190

170

150

130

no

900*F IOOO«F IIOO^F I200®F

PRECIPITATION HARDENING TEMPERATURE

Figure 15. Energy absorbed versus precipitation hardening

temperature for A710 austenitized at 1650°F

for 90 minutes and then aged at the indicated

temperature and time.

42

JOULES
* m.

3*ooa

Ogs

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B
500«C 550«C 600®C 650«C

200

270

250

230

210

190

170

150

110

PRECIPITATION HARDENING TEMPERATURE

Figure 26. Energy absorbed versus precipitation hardening

temperature for A710 austenitized at 1650°F

for 30 minutes. Letters C and D refer to the

type of fracture exhibited on the broken impact

specimens. C is cleavage, and D, ductile

fracture.

53

JOULES
^^•">005/ '?*O0‘i i"*ooo< ^•ooe

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500®C

“1

550«C 600«C eso^c

PRECIPITATION HARDENING TEMPERATURE

Figure 27. Energy absorbed versus precipitation hardening

temperature for A710 austenitized at 1650°F

for 90 minutes. Letters C and D refer to the

type of fracture exhibited on the broken impact

specimens. C is cleavage and D, ductile

fracture.

54

JOULES
J. -‘'HI' 'I'

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JOULES

100 130 160 190 220 250 280 310

ENERGY ABSOR BED , FT - LB

Figure 28. Ultimate tensile stress and lower yield stress versus energy

absorbed.

55

MPa
WbO

OSS OOS Q8i ,^031 OW ^OSI 001

V aj “ Tt , a38i^08SA YcW3&13

iili- vl^fmd 2ea*(^£' t«»s $vj-«mrjru »8S yiuc^^l
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750

700

650

600

550

500

Figure 29. Ultimate tensile stress and lower yield stress versus hardness,

HRB.

56

MPa

HARDNESS,

HRB

JOULES

100 130 160 190 220 250 280 310

Figure 30. Hardness, HRB, versus energy absorbed for each test condition.

57

500®C 550«C 600«C 650«C

PRECIPITATION HARDENING TEMPERATURE

Figure 31. Upper yield stress minus lower yield stress versus

precipitation hardening temperature for A710 austenitized at

1650°F for 30 minutes and then aged at the indicated

temperature and time.

58

MPa
!| <

5

r ^

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500®C 550®C 600®C 650®C

60

50

40

30

20

10

PRECIPITATION HARDENING TEMPERATURE

Figure 32. Upper yield stress minus lower yield stress versus

precipitation hardening temperature for A710 austenitized at

1650°F for 90 minutes and then aged at the indicated

temperature and time.

59

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500«C 550«C 600®C 650«C

PRECIPITATION HARDENING TEMPERATURE

Figure 33. Ratio of upper yield stress, lower yield stress to the

ultimate tensile stress versus precipitation hardening

temperature for A710 austenitized at 1650°F for 30 minutes and

then aged at the indicated temperature and time.

60
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1.00

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Figure 34.

500°C 550®C 600®C 650°C

UYS/UTS RATIO I650**F-90MIN- W/Q —-- LYS /UTS RATIO • 30MIN - PH - A/C

60MIN-PH-A/C

A 90MIN-PH-A/C

900®F 1000® F IIOO®F I200®F

PRECIPITATION HARDENING TEMPERATURE

Ratio of upper yield stress, lower yield stress to the ultimate

tensile stress versus precipitation hardening temperature for

A710 austenitized at 1650°F for 90 minutes and then aged at the

indicated temperature and time.

61
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Acknowledgment

The authors thank L.K. Ives for his efforts in the electron microscopy

investigation. And we are pleased to acknowledge our secretaries, Ms. D.

Easterday and Mrs. P. Salpino, both of whom worked so diligently on the

manuscript.

63
APPENDIX A.

TITLE:

SUBMITTED

Small Angle Neutron Scattering (SANS)

Study of Aging of ASTM A710 Grade A, Class 3

Steel

TO: Dr. Richard Fields and George Hicho

Fracture and Deformation Division

National Bureau of Standards

Gaithersburg, MD

BY: Dr. Satish P. Singhal
INTRODUCTION

Recently, a low carbon precipitation hardenable alloy steel, A710, has

been considered for high strength, welded structures. Strengthening of A710

takes place by controlled precipitation of copper rich particles in a ferrite

matrix at temperatures below 1300°F. The goal of this small-angle neutron

scattering (SANS) study was to find a correlation between aging treatment

(duration and temperature), mechanical properties (hardness), and SANS

characteristics.

SAMPLES AND EXPERIMENTAL PARAMETERS

Prior to aging all samples ( 0.11m thick) were austenitized at 1650°F for

about 90 minutes. Solution treatment was terminated by water quenching.

Aging was carried out in the temperature range of 900 to 1200°F for times 30

min to 90 min. Other sample details are given in Table 1. SANS

investigations were carried out in the scattering vector (Q) range 0.001 to

0.013 nm”^. In this Q range the contribution of the precipitates in the size

range of 4.8 to 60nm is measured.

The detailed theory for data reduction has been discussed by Kostorz (2).

A few salient features, however, are described here. Let us assume that Ip is

the intensity scattered by the sample containing the precipitates, then

following Kostorz (2):

65
Ir^(Q) “ ^ p ^0 p p

dzP

SCatt

dfi (1)

where

M = total monitor count,

P

(f)^ = neutron flux on the sample - per monitor count,

t = sample thickness,

P T = sample transmission

P = detector efficiency

A = solid angle of one detector element at the scattering

vector Q,

^^Latt = macroscopic differential scattering cross-section of the sample

Differential scattering cross-section is the "bulk averaged" measure of

the precipitate volume fraction, precipitate size distribution, and

"scattering contrast" between the matrix and precipitates. Scattering

contrast in this case depends upon the difference in the Cu content in the

matrix and the precipitates. Additional contrast comes into the play due to

the magnetic nature of the ferrite matrix as precipitates themselves are

non-magnetic.

It is clear from equation 1 that as long as we divide all the sample's

intensity (Ip(Q)) by the thickness (tp) and transmission (Tp)

we can intercompare the spectra from various samples after adjusting the

counts to the same monitor value (Mp). 4>o and ed are the instrumental

constants and do not change, whereas the program "AVERAGE" normalizes for

the variation of with Q, as well as calculates the intensities at each

scattering vector. Further details of data processing will be described

in the results and discussion section. Other SANS experimental parameters

are given in Table (2)

.

66
CONSIDERATIONS RELATING SANS WITH PRECIPITATION PHENOMENA.

The central idea in understanding the precipitation phenomena in

supersaturated solid-solutions is that first nucleation takes place and then

these nuclei grow by diffusion controlled mechanisms. For an alloy of fixed

composition the aging temperature dictates the number of particles (nuclei)

_3 cm in the specimen and aging time determines their size distribution. At

_3

longer aging times, the number of particles cm (N^) reduces because large

particles grow at the expense of smaller ones by a mechanism called "Ostwald

ripening". Following Guinier (2) for a system of "randomly" distributed

particles the macroscopic cross-section (de/dfi) (see eq. 1) can be given as:

Here

R

=

• ''n• ex p(-Q^-R^ ) +*(Q) (2).

Ill p 2

Scattering length density of the matrix (m) and precipitate

(P).

Volume of one precipitate (V ) and number of precipitates in

the sample (Np). P

Radius of gyration related to the particle diameter as

D = 2R /5/3.

i|i(Q) = Interparticle interference function which is negligible in case of low volume fractions of precipitates in the matrix.

In cases where (Q) is significant a peak at is seen in the scattering

curve.

The Guinier law, Eq 2, has some limitation in the sense that it considers

the particulate assembly to be monodispersoid. One could, however, analyze

low Q regions of the scattering curves in order to get useful information

about the "average" size growth of the precipitates with thermal treatment.

One can also see from Eq 2 that with an increase of R (or D), the slope of a

67
2 log dz/dr vs Q curve will increase, signifying the size growth of the

precipitates. Where such growth takes place, the scattering curve translates

to lower scattering vectors, i.e., intensity at low Q rises but at high Q, it

goes down. This is schematically shown in figure 35. The temporal cross-over

between two subsequent scattering curves (e.q. at Qc. in Figure 1) is the

hallmark of such coarsening process.

RESULTS AND DISCUSSIONS

General Remarks

Since precipitation phenomena in A710 involves the nucleation and growth

of a non-magnetic phase within the magnetic ferrite matrix, it is necessary to

eliminate magnetic scattering from the matrix in order to measure only the

nuclear (contrast) scattering between the precipitate and matrix. This was

accomplished by applying a 6K Gauss magnetic field in the direction

perpendicular to the neutron beam across the sample. The resultant

2-dimensional detector pattern for the sample #23 is shown in Figure 2. One

can see that the resulting intensity in the field direction (H) decreases

because all the moments are aligned along it. This results in a familiar

"figure of eight" pattern seen very commonly for magnetic samples. In Figure

2, direction H, one sees only nuclear contrast scattering whereas in the

direction perpendicular to H, one sees nuclear plus magnetic scattering.

Similar patterns were seen for all other samples as well. In Figure 3 we show

the scattering curves in the field direction and perpendicular to it. This

averaging was done by cutting a 30° sector along the H and perpendicular to

the H (Figure 2) and averaging both sides of the detector. One can see

(Figure 3) that for a given Q the scattered intensity in the field direction

is always lower than in the perpendicular direction.

68
Some interesting conclusions can be drawn from the Guinier plot in Figure

3. In a monodispersed system, where all the precipitates have the same size,

this plot should be a straight line. This is obviously not the case here as

2 one can fit a straight line only to the very narrow portion of the log I vs Q

curve (Figure 3) in the low Q region. (This is further el lustrated by Figure

7). Also, the absence of any peaks in the scattering curve indicates that

particle density is low enough to neglect the interparticle interference T (Q)

in Eq. 2. In the absence of ideal behavior, complex deconvolution schemes are

required for obtaining the size distribution of the particles. Some attempts

in this direction are described in the literature.

SANS characteristics of aged samples are shown in Figure 4. According to

the theory of nucleation and growth, for an alloy of asymmetric composition,

as one increases the temperature of aging the equilibrium volume fraction of

the second phase decreases because the miscibility gap becomes narrower at

high temperatures. However, as long as the reaction is controlled mainly by

the interdiffusion coefficient, the volume fraction precipitated in the "same"

time period will increase as the aging temperature increases*. This is very

nicely indicated by Figure 4. Overall integrated intensity is lowest for #21

as this curve lies below the rest in the scattering pattern (Figure 4). Upon

aging at 1000° F; there is a significant rise in the

* This behavior will be obeyed until the temperature is high enough to

intersect the nose of the 'C curve in the TTT diagram, above which

temperature the reaction is driving force controlled instead of diffusion

control led.

69
intensity in the Q region 0.003 to O.OOSnm"^, signifying an increase in the

volume fraction of the precipitates in the size range 80 to 20.0nm.

With a further increase in the aging temperature to 1100°F, a cross-over

appears between the scattering curves which move to lower Q values with

continuing increases in the temperature (see specimens 22, 23, and 24). This

signifies that due to more rapid diffusion at higher temperatures, enough

"growth" and coarsening is taking place in 90 minutes that the size

distribution is shifted to larger sizes and hence the scattering curve shifts

to lower Q values.

Correlation of the aforementioned results with the hardness data is not

direct. From Table 1, we can see that hardness is highest for #21 and it

decreases as one ages isochronal ly at higher temperatures. According to the

prevalent theories of precipitation hardening, there are two dislocation

mechanisms which can control the strength of an age-hardened material. These

are (1) dislocation cutting, and (2) dislocation bowing. If precipitates are

softer than the matrix, then dislocations can easily cut through the second

phase particles. It is therefore desirable that precipitate elastic constants

be higher than the matrix. The bowing of dislocations can be prevented by

obtaining a fine dispersion of the overall volume fraction of precipitates

such that the inter-particle distance is minimum. If the inter-particle

distance is low, then a higher shear stress is required to bow a dislocation

through the precipitates. This action results in higher strength. Even though

it is clear from Figure 4 that the volume fraction of copper precipitates

present in sample #21 is lower than in the other samples, the distribution is

finer. This results in a smaller interparticle spacing compared to the other

samples. Therefore, a higher hardness is obtained. One conclusion which

70
becomes clear immediately in this case is that as hardness decreases (from #21

to #24), the scattered intensity in the high Q region increases, but in the

low Q region, it decreases. This observation also points to a shift of

precipitate size distribution to larger sizes.

Another interesting correlation is revealed when we plot total neutron

cross-section (Table 1) versus the hardness values (Figure 5). As hardness

increases, the total neutron cross-section decreases. This decrease is

equivalent to the sample becoming more transparent to the neutron beam (or the

sample transmission per unit thickness increasing). This observation is

consistent with Figure 4 in the sense that the scattering curve for the sample

with lowest total neutron cross-section lies underneath all the other curves

(#21 in Figure 4). Such total cross-section increase with increasing aging

temperature is indicative of the fact that the overall integrated SANS

intensity, and therefore the volume fraction of the second phase, increases

with an increase in aging temperature. The simplicity of the correlation

between hardness and the total cross-section for the isochronally aged samples

is rather striking and thus has some potential for being a good NDE tool.

Only sample numbers 14 and 22 provided for this study are from an

isothermal ly aged lot. According to the theory of nucleation and growth and

its correlation to SANS (see Figure 1), a translation of SANS curves with

aging time is expected during the isothermal evolution. Figure 6 shows

scattering curves for the A710 samples aged for 30 and 90 minutes respectively

at 1000°F. We can see that due to the additional hour of aging, the scattered

intensity increases over almost the entire Q range, except that a slight

decrease is observed in the small portion of the high Q region. In

71
this case the hardness and total cross-section trend similar to the

isochronally aged samples is followed. Sample #22, which is a strong

scatterer compared to #14, also has a larger total cross-section, but a lower

hardness (98.5 HRB vs 99.8 HRB) (see Table 1).

We have interpreted the translation of the SANS curves with the increase

in aging temperature (figure 4) as a shift of precipitate size distribution to

the higher sizes. In the absence of detailed size distribution analysis, this

fact should also be reflected by the increase in Guinier radius with the aging

temperature. Hence a "Guinier" analysis was carried out for the spectra of

all samples. NBS program DATAPLOT was used for fitting Eq. (2) (with T (Q)=0)

to the relevant part of the SANS curves. An example of this fit is shown in

Figure 7, with the blank squares indicating the data points and a continuous

line as the fitted line. For sample #14, the radius of gyration or "Guinier"

radius was 17.7nm, which translates to a particle size of 45.7nm. Radius of

gyration for other samples is given by the last column of Table 1. It is our

belief that all the Rg values marked by an asterik are unreliable. Let us

look at Figure 4 and samples 21 to 24. There, the scattering curve translates

towards the main beam. Under such a situation, not only do the resolution

corrections become progressively large, but the real "Guinier region" also

becomes inaccessible. Therefore in order to carry out appropriate Guinier

analyses for samples 22, 23, 24, further measurements must be performed in the

Q region much lower than the lowest (O.OOlnm”^) covered in this study. This

also suggests that samples heat treated at temperatures above 900°F have a

large fraction of particles larger than 60nm in size.

72
Table 1. Test sample details

Sample Aging Hardness Thickness Transmission Total Neutron Guinier

Number Treatment Cm at X =6.2A Cross Section Radius

(cm RG(nm)

30min/1000°F 99.8Rg , 0.293 0.76139

90min/900°F 23.2Rp

-100. 2RB

0.291 0.77241

90min/1000°F 98.5Rg 0.2909 0.75254

90min/1100°F 93.8Rg 0.2921 0.74691

90min/1200°F 89.0Rg 0.2995 0.73738

*Reliability of these Rp values is questionable,

explanation.

0.93041 17.707

0.88742 17.97

0.97732 16.64*

0.99901 15.7*

1.0172 16.44*

See the text for relevant

73
Table 2. SANS parameters for A710 Steel

Wavelength - 6.2 A° ± (ax/X'0.25)

Monochromator - Mechanical Velocity Selector

Source of Neutrons - 10 MW Reactor

Filters - 15" Be and 10" Bi Single Crystal

Detector Size - 65X65 cm 110X110 active channels

Detector gas - ^He + Xe

Detector distance - 3.6 meters

Beam Size - 1.2 cm.

Col 1 imation - iris 27mm - 12mm, 4.1 meters apart

Sample to entrance collimator distance 4.50 meters.

74
INTENSITY

SCATTERING VECTOR

Figure 1. SANS results showing intensity versus scattering

vector, t^ and t„ denote aging times at a fixed

temperature.

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76
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LOG

INTENSITY

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Figure 3. Guinier plot of log intensity versus the

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NBS>n4A (REV. 2-80

U.S. DEPT. OF COMM.

BIBLIOGRAPHIC DATA

SHEET (See instructions)

1 . PUBLICATION OR

REPORT NO.

NBSIR 83-2685

2. Performing Organ. Report No, 3. PubI ication Date

May 1983

4. TITLE AND SUBTITLE

Effects of Variations in Heat Treatment on the Mechanical Properties and

Microstructure of ASTM A710 Grade A Class 3 Steel

5. AUTHOR(S)

G. E. Hicho, L. C. Smith, S. Singhal, and R. J. Fields

6. PERFORMING ORGANIZATION (If joint or other than MBS. see instructions) 7. Contract/Grant No. national bureau of standards

DEPARTMENT OF COMMERCE 8. Type of Report & Period Covered

WASHINGTON, D.C. 20234

NBSIR

9. SPONSORING ORGANIZATION NAME AND COMPLETE ADDRESS (Street. City. State. ZIP)-

David Taylor Ship R§D Center

Annapolis, MD

10. SUPPLEMENTARY NOTES

I I Document describes a computer program; SF-185, FIPS Software Summary, is attached.

11. A,BSTRACT (A 200-word or less factual summary of most si gn i fi cant information. If document includes a significant

bibliography or literature survey, mention it here)

A final report assessing the effects of variations in heat treatment on the mechanical

properties and microstructure of ASTM A710 Grade A Class 3 steel is given. The goal

of this project was to evaluate the sensitivity of the mechanical properties of A710,

an age hardening steel, to heat treatment variations that may occur in production

of thick plate. Tensile and impact tests were conducted on the material heat

treated under the various conditions. Microstructure evaluations, using optical

and electron metallographic methods were conducted. Small angle neutron scattering

examinations were also performed on selected specimens in order to evaluate the

precipitate distribution. Results indicate that austenitization and aging time

has little effect on the resultant mechanical properties. However, the mechanical

properties were found to be sensitive to aging temperature. Impact results of

0 F indicate that aging at 900 F is detrimental, in that cleavage was found on the

fracture surface of the broken impacts. Furthermore, the results obtained here

suggest that welding or similar heating operations will not seriously degrade this

material's properties.

12. KEY WORDS (Six to twelve entries; alphabetical order; capitalize only proper names; and separate key words by semicolons)

Cleavage fracture; ductile fracture; energy absorbed; lower yield strength;

microstructure; precipitate; ultimate tensile strength; upper yield strength.

13. AVAILABILITY

1 1 Unlimited

y 1 For Official Distribution. Do Not Release to NTIS

1 1 Order From Superintendent of Documents, U.S. Government Printing Office, Washington, D.C. 20402.

! 1 Order From National Technical Information Service (NTIS), Springfield, VA. 22161

14, NO. OF

PRINTED PAGES

15. Price

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