NBSIR 80-1627

Reference NBS

Publl - cations

TECHNICAL REPORTS

MATERIALS STUDIES FOR

MAGNETIC FUSION ENERGY

APPLICATIONS AT LOW TEMPERATURES

j I I TEMPERATURE

STRESS INTENSITY RANGE

TO

DEPARTMENT OF ENERGY

OFFICE OF FUSION ENERGY

WASHINGTON, D.C. 20545

-QC

100

.U56

80-1627

1980

FRACTURE AND DEFORMATION DIVISION

NATIONAL BUREAU OF STANDARDS

BOULDER, CO 80303

NBSm 80-1627

ATIOIIAL BUREAU

or STANBABM

DBRAKT UC 20 b,c

SEP / 9 1980

TECHNICAL REPORTS Ko ' , MATERIALS STUDIES FOR

MAGNETIC FUSION ENERGY

APPLICATIONS AT LOW TEMPERATURES - III

Edited By

R.P. Reed

Fracture and Deformation Division

National Bureau of Standards

Boulder, CD 80303

June 1980

Sponsored By

Department of Energy

Office of Fusion Energy

Washington, D.C. 20545

U S. DEPARTMENT OF COMMERCE, Philip M. Klutznick, Secretary

Luther H. Hodges, Jr., Deputy Secretary

Jordan J. Baruch, Assistant Secretary for Productivity, Technology and Innovation

NATIONAL BUREAU OF STANDARDS, Ernest Ambler, Director
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CONTENTS

Page

SUMMARY 3

INTRODUCTION 5

HIGHLIGHTS OF RESULTS 7

STRUCTURAL ALLOYS 13

Interstitial Carbon and Nitrogen Effects on the Tensile and

Fracture Parameters of AISI 304 Stainless Steels 15

Temperature Dependence of Yielding in Austenitic

Stainless Steels 49

Anomalous Low- Temperature Elastic-Constant Behavior in

Fe-13Cr-19Mn 79

Room-temperature Elastic Constants and Low-Temperature Sound

Velocities for Six Nitrogen-Strengthened Stainless Steels 91 Sound Velocities and Elastic Constants

of Austenitic Steels 304, 310, and 316 105

Elastic Constants of Stainless Steels 304, 310, 31 6

between 4 and 295 K 117

Young's Modulus of a Copper-Stabilized

Niobium-Titanium Superconductive Wire 133

WELDING 151

Weldments for Liquid Helium Service 153

Fracture Behavior of Ferrite-Free Stainless Steel Welds

i n Li qui d Hel i urn 1 67

Weld Process Study for 31 6L Stainless Steel Weld Metal

at Cryogenic Temperatures 195

The Computer-Aided J-Integral Test Facility at NBS 205

Tensile and Notch Tensile Properties at Room Temperature

and 4 K of Welds in Aluminum Alloy 5083-0 Plate 217

Note on a Simple Test Method for Estimating 245

NONMETALLICS 257

Nonmetal lies for Magnet Systems 259

Cryogenic Properties of Fil amentary-Reinforced Composites:

An Update 263

Mechanical Properties of G-IOCR and G-llCR 333

Mechanical, Electrical, and Thermal Characterization of

G-IOCR and G-llCR Glass/Cloth Epoxy Laminates Between

Room Temperature and 4 K 341

Temperature Dependence of Young's Modulus and Internal

Friction of G-IOCR and G-llCR Epoxy Resins 367

Directional Variation of Young's Modulus and Internal

Friction in Tv/o Fiberglass-Cloth-Epoxy Composites 379

Thermal Conductivity of G-IOCR and G-llCR Industrial

Laminates at Low-Temperatures (Part II) 395

Thermal Expansion of Cryogenic-Grade Glass-Epoxy Laminates 405

TECHNOLOGY TRANSFER 413

Structures, Insulators and Conductors for Large

Superconducting Magnets 415

VAIL WORKSHOP 425

PROGRAM DESCRIPTION 433

ORGANIZATIONAL CONTACTS

1

SUMMARY

The reports presented here summarize the work of the low temperature

materials research project for the third year of the program. The various

projects are outlined and the research results are presented. The major

thrust of the measurements has been the evaluation of the low temperature

mechanical and physical properties of stainless steel base metal and welds,

with particular emphasis on the nitrogen-strengthened stainless steels.

Aluminum alloys have receive'd some consideration also. Work has also

been done on the production and standardization of nonmetal 1 ics , primarily in

dustrial laminates, for low temperature applications and on the measurement

of their properties at cryogenic temperatures . A brief description of

the third NBS/DoE Vail workshop, held in October 1979, is given.

Note

Certain commercial equipment, instruments, or materials are identified in this paper to adequately specify the experimental procedure. In

no case does such identification imply recommendation or endorsement

by the National Bureau of Standards, nor does it imply that the material

or equipment identified is necessarily the best available for the purpose.

Papers by non-NBS authors have not been reviewed or edited by NBS. Therefore

the National Bureau of Standards accepts no responsibility for comments

or recommendations contained therein.

INTRODUCTION

This report contains results of a research program to produce

materials properties data to facilitate design and development of cryogenic structures for superconducting magnets required for plasma confinement in magnetic fusion energy power plants and prototypes. The

program was conceived and developed jointly by the staffs of the National

Bureau of Standards and the Office of Fusion Energy of the Department of

Energy. This program, sponsored by DoE, is managed by NBS. Research is

conducted at NBS and at various other laboratories through subcontracts

with NBS.
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Highlights of Results

7

HIGHLIGHTS OF RESULTS

We present here a brief description of outstanding results from the

project reports, which follow in the next section.

Alloys for Low Temperature Service

The effects on 4 K strength and toughness of carbon and nitrogen additions to an AISI 304L base alloy were examined. Amounts of carbon between

0.03 and 0.09 weight percent, and of nitrogen between 0.03 and 0.22 weight

percent were used. Three- fold increases of yield strength were obtained

at 4 K. The strength was found proportional to carbon plus nitrogen additions. The dependence of flow strength on carbon and nitrogen concentration is more pronounced at lower temperatures . Although toughness decreases

and fatigue crack growth rates increase with carbon and nitrogen additions,

adequate toughness and crack growth rates are maintained at 4 K in the

alloys with high concentrations.

The temperature dependence of the flow strength of AISI 316 was

measured at about ten different temperatures between 4 and 300 K. Unlike

AISI 304L, no anomalous decreases of flow strength at lower temperatures

were observed. Using the Oak Ridge National Laboratory high temperature

data for the same alloy heat, equations predicting the temperature dependence of alloys were assessed. The best low temperature fit to our

data relates flow stress to the reciprocal of temperature. The relationship of this temperature dependence to physical factors affecting strength

is being examined.

Physical Properties of Stainless Steels

This report contains brief descriptions of four experimental studies

on the elastic properties of 300-series austenitic stainless steels:

9
1. Room-temperature sound velocities and elastic constants of

austenitic stainless steels 304, 310, and 316. This study extends our

previously reported effort on the elastic-constant variability in stainless steel 304. For all three steels, for all elastic constants, variabilities are less than two percent.

2. Elastic constants of stainless steels 304, 310, and 316 between

4 and 295 K. All three steels show low-temperature elastic-constant

anomalies related to magnetic transitions. Stainless-steel 310 behaves

differently in this regard than 304 and 316.

3. Room-temperature elastic constants and low-temperature sound

velocities in six nitrogen-strengthened austenitic stainless steels.

Despite large composition differences, room-temperature elastic constants

vary only slightly in these alloys. At low temperatures , elastic constants, affected by magnetic transitions, show qualitatively similar,

but quantitatively different, behavior.

4. Anomalous low-temperature elastic-constant behavior in Fe-13Cr19Mn. This Soviet stainless steel, containing high manganese, but no

nickel, exhibits large elastic-constant anomalies upon cooling to about

290 K.

Wei di ng

Tensile and fracture toughness tests were conducted on stainless

steel welds at 4 K in support of the Large Coil Project (LCP) and Magnetic Fusion

Test Facility programs. Welds were provided by LCP contractors and Lawrence

Livermore Laboratory (LLL). Notable results at 4 K are summarized below. Peering the weld after each pass improved the toughness of the LCP

316L welds. Annealing the LCP 316L welds resulted in a substantial improvement in toughness. The 316L SMA welds supplied by LLL had satisfactory

10
strength and toughness at 4 K. An open-arc flux-cored 308L weld exhibited

brittle behavior in the 4 K toughness tests; this was attributed to

nitrogen pickup during welding. Experimental weld compositions supplied

by filler metal producers were evaluated at 4 K, The 1 8Cr-l 6Ni -9Mn

alloy had a yield strength of 1030 MPa and a toughness of 121 MPa

at 4 K. The 1 6Cr-l 6Ni -2Mo alloy deposited by the FCAW process had a

toughness of 239 MPa/iri and a yield strength of 607 MPa at 4 K.

Alcoa conducted an evaluation of 51 mm thick 5083-0 aluminum weldments provided by several fabricators. All welds were made using 5183

electrodes and the gas metal-arc process. The tensile and notch tensile

properties were determined at room temperature and 4 K. The same test

plates will be fracture toughness tested by NBS in FY80 to extend to

4 K the correlation between notch-yield ratio and fracture toughness,

which is currently valid at room temperature and 77 K.

Teledyne-McKay prepared a series of six test weldments on AISI 304

stainless steel base plate using three welding processes: GTA, GMA, and

SAW. Each of the weldments was tested to measure the tensile properties

at room temperature, the Charpy impact toughness at 77 K, and the microfissuring tendencies. NBS will evaluate the. tensile properties ana

fracture toughness at 4 K during FY80. The Charpy values indicate that

the GTA and GMA welds meet ASME requirements at 77 K, but the SA welds

are brittle.

Nonmetal lies for Magnet Structures

The initial static mechanical property characterization of the

G-IOCR and G-llCR cryogenic grades of industrial glass-epoxy laminates

have been completed. Additionally, the CR-grade products produced by a

11
second manufacturer have been evaluated and the performance of these

grades fabricated with boron-free type E glass fabric has been assessed.

No statistically significant differences in cryogenic tensile or compression strength or in tensile modulus was observed in the tested

products. The CR grades have proven to have excellent low temperature

properties and have been accepted by a variety of fabricators for use in

superconducting magnet technology.

Material representing all CR-grade production has been sent to ORNL

for assessment of neutron radiation performance at 76 K.

NBS has worked closely with the Owens-Corni ng Fiberglas Corporation

to systematize the available data on glass-mat reinforced epoxy laminates

for use at cryogenic temperatures. These materials provide an alternative

to fabric-reinforced products where somewhat lower mechanical properties

can be tolerated at a substantially lower laminate cost.

Assessment of the magnitudes and sources of variability in cryogenic properties of the CR-grade laminates will continue in FY80. We

shall begin a program to assess the level of variability and the lower

bounds on cryogenic properties of cotton fabric/phenolic products of

NEMA L and LE types. We shall begin a literature search to define the

present state of knowledge on the cryogenic performance of nonmetallic

insulating films of interest to the MFE community.

Technology Transfer

The annual NBS/DoE Workshop on Materials at Low Temperatures was

attended by about seventy representatives of the materials and magnet

design communities at Vail, Colorado in October 1979. Again, the meeting

was viewed constructively; it served as a focal point to identify and

discuss magnet construction or design problems related to materials.

For a few who have become recently involved in cryogenic applications,

the workshop served an educational role.

12
Structural Alloys

13

INTERSTITIAL CARBON AMD NITROGEN EFFECTS ON THE

TENSILE AND FRACTURE PARAMETERS OF AISI 304

STAINLESS STEELS

National Bureau of Standards

15

INTERSTITIAL CARBON AND NITROGEN EFFECTS ON THE

TENSILE AND FRACTURE PARAMETERS OF AISI 304

STAINLESS STEELS

R. L. Tobler and R. P. Reed

National Bureau of Standards

Boulder, Colorado

17
ABSTRACT

This study was conducted to determine the effects of carbon and

nitrogen on the mechanical properties of an Fe-1 8Cr-l Ofli austenitic

stainless steel having a base composition corresponding to AISI 304.

Nine steels having carbon plus nitrogen (C+N) contents between 0.067 and

0.325 wt.% were tensile tested at 295, 195, 76, and 4 K. J-integral

fracture toughness and fatigue crack growth rate tests using 25-mm

thick, compact specimens (TL orientation) were also performed at 4 K.

At 295 K, the yield stress was linearly dependent on C+N content, with

values ranging from 222 to 358 MPa. At cryogenic temperatures , the

interstitial strengthening effects were magnified, so that the yield

stress values ranged from 329 to 1177 MPa. The fatigue crack growth

resistance and fracture toughness decreased with increasing C+N content.

The estimates obtained at 4 K ranged from 337 to 123 MPa-m^^^,

exhibiting an inverse dependence on yield. Measurements of ultimate

stress, elongation, and reduction of area are reported, and the role of

martensite transformation in these steels is discussed.

INTRODUCTION

The AISI 300 series stainless steels are austenitic Fe-Cr-Ni

alloys, offering relatively low strength but excellent cryogenic ductility and toughness^^ ^ • By virtue of favorable mechanical properties,

service history, and availability, AISI 304 is perhaps the most widely

used cryogenic alloy in the world. But ni trogen-strengthened grades,

such as AISI 304 N or AISI 304 LN,have attracted attention as possible

substitutes for AISI 304 in applications demanding higher strength.

18
Nitrogen is a relatively inexpensive and effective strengthener

Nitrogen also stabilizes the austenitic structure and reduces the probability of martensite transformation . Carbon has similar effects, but

carbon contents in austenitic stainless steels must be held to low

levels to prevent sensitization during thermal excursions.

Ni trogen-strengthened austenitic stainless steels are gaining

popularity, but some applications are hindered by limited availability,

limited service experience, lower fracture toughness, and a lack of

[ 2 ] cryogenic design data . Several mechanical property studies were conducted on these steel s^^~^

but most papers are concerned only with

room temperature behavior. The present study was undertaken to describe

cryogenic behavior in greater detail. Yield stress, ultimate stress,

elongation, and reduction of area measurements for nine heats of an

Fe-18Cr-10Ni stainless steel are reported at 295, 195, 76, and 4 K.

The carbon and nitrogen contents of the nine heats were varied to

enable a systematic study of interstitial concentration on mechanical

property trends. The results can be used to predict the properties of

commercial alloys of known compositions or to tailor the properties for

a particular application by appropriate specification of the steel

composition.

MATERIALS

Nine stainless steel plates were purchased from a commercial

steel manufacturer. These plates had a nominal base composition

falling within the limits set by ASTM specification A240 for AISI 304

stai nl ess steel : Fe-18Cr-l ONi -1 . 5Mn-0. 02P-0. 02S-0. 55Si-0.2Mo-0.2Cu

The carbon and nitrogen levels varied however, with carbon at 0.03,

19
O.OG, or 0.09 wt.%, and nitrogen at 0.04, 0.12, 0.24 wt.%. The mill

chemical analyses are listed in Table I. The nine plates were produced

from three 136-kg (300-lb) vacuum-induction-melted heats, split with

respect to carbon level, and teemed into 76-cm x 200 cm x 360-cm hottopped cast-iron ingot molds. The ingots were then reheated and soaked

at 1561 K (2350°F), hot-rolled to 25.4-mm thick plates, and air cooled.

The plates were finally annealed at 1332 ± 7 K (1937 ± 13°F) for 1 h

and water quenched. The hardness and grain size measurements are listed

in Table I.

EXPERIMENTAL PROCEDURE

Tensi 1

Tensile tests were conducted using the cryostat described by

Reed^^^^. Low temperatures were achieved by immersing the specimen and

cryostat in alcohol and dry ice at 195 K, liquid nitrogen at 76 K, and

liquid helium at 4 K. The specimen geometry is shown in in Fig. 1.

The load-vs-deflection curves were recorded using the outputs from a

commercial load cell and strain gage extensometer . The specimen yield

stress at 0.2% plastic strain, ultimate tensile stress, elongation, and

reduction of area were measured.

Fatigue Crack Growth Rates

Fatigue and fracture toughness tests were performed using the

r 1 oi

lOO-kN machine and cryostat described by Fowlkes and Tobler*- These

tests used compact specimens in the TL orentation, as defined by ASTM

n 91 Method E-399-74 . The specimen geometry is shown in Fig. 1; specimen

width, w, was 50.8 mm, and thickness, B, was 24.5 mm. These specimens

20
were precracked at 77 K, using a maximum load of 23 k.N. Fatigue crack

growth rates were then measured at 4 K. The load ratio was 0.1 and

the load cycle was sinusoidal at a frequency of 25 Hz. The compliance

method of crack growth rate _measurement was used, as described pren 8l viously*- . Data were obtained from one or two specimens per heat of

steel

Fracture Toughness

J-integral tests were necessary because these austenitic stainless

steels fail in a ductile (elastic-plastic) mode with gradual crack

extension at 4 K. Therefore, single-specimen J-integral tests were used

to measure fracture toughness. The technique was a variant of the unr 2 201 loading compliance method*- ’ modified by incorporating a computer to

enable digital data acquisition and J-resistance curve plotting during

the tests.

The fracture specimens were fatigue cracked at 4 K to a crack

length, a, between 28.8 and 36 mm (g 0.57 to 0.71). The specimens

were then loaded in displacement control at the rate of about 50 mm/s.

As the load increased, partial unloadings were performed periodical ly.

The load reduction during unloading was never greater than 10% of the

maximum load. Using the crack-length-vs-compl iance correlation, the

crack length at each unloading was inferred and used to obtain the

crack extension increment (Aa) from the difference of the initial nrecrack

length. The J value at each unloading point was calculated using expressi
T21

where A is the Merkl e-Corten factor*- -*

A is the area under the load-vs-defection curve at the unloading

point, and

b is the specimen ligament (b = W - a).

J-resistance curves consisting of many sets of J and Aa points were

generated, as shown by the example in Fig. 2.

The Jx value for each steel was taken at the intersection of the

Ic

J-resistance curve with the blunting line. The blunting line is given

by J = 2a^ A^ where is the flow stress, which is equal to the average

of the yield and ultimate tensile stresses. As shown in Fig. 2, there

is an artificial offset of the data from the theoretical blunting line,

so a blunting line parallel to the theoretical blunting line but passing

through the data was used. After the test, the Aa values inferred from

the compliance data agreed within 10% with those found by direct measurement from the fracture surface. Scatter in the J-resistance curves lead

to uncertainties of ±12% in the values of Jj^.

The validity of Jj^ results was determined by the size criterion

[221 which requires that*-

B, b, a ^ Ic

Of

Previous results have shown that valid results at 295 K are impossible

for annealed austenitic stainless steels of 24.5-mm thickness, so a

complete temperature dependence of Jj^, was not attempted. Only 4 K

values are reported.

22
Phase Transformation

In all low temperature tensile, fatigue, and fracture tests,

partial transformation of the paramagnetic austenite to ferromagnetic

a' (bcc) martensite was observed. The extent of a' transformation was

[231 estimated using a bar-magnet device described previously . Small

amounts of e (hep) martensite may also have formed, but this was not

measured.

RESULTS AND DISCUSSION

The tensile properties of the steels tested in this study are

listed in Table II and plotted in Figs. 3-6. As shown in Fig. 3, the

yield stress at 0.2% plastic strain is a linear function of interstitial

C+N content at 295 K, where the yield stresses vary from 222 to 358 MPa.

With greater data scatter, linear dependences are also observed at cryogenic temperatures , but the effectiveness of interstitial strengthening

progressively increases as temperature is reduced. Thus, the 4 K yield

stresses have greatly increased, to values ranging from 329 to 1177 MPa.

As shown in Fig. 3, there are departures from the linear trends at

low temperatures for C+N contents of 0.128 wt.% or less. These breaks

from otherwise linear dependences are attributed to deformation mode

changes, possibly in association with austenite instability. It is

known that partial phase transformations to a' (bcc) or c (hep) martensite phases can occur in AISI 304 alloys having low interstitial contents

Such transformations are not probable at 295 K, but low temperatures

tend to promote transformations . An x-ray examination of selected

specimens after only 0.2% plastic strain would be necessary to test this

interpretation

23
In Fig. 4 the yield stress daita are plotted versus temperature

using a semilog scale. Previous work showed that many stable austenitic

stainless steels exhibit a linear trend when their a values are plotted

r?5l ^

in this way"- , whereas less stable alloys, such as AISI 304L, exhibit

less regular behavior. The AISI 394L behavior is marked by anomalous

F41 decreases of strength in the temperature range between 200 and 100 K .

The trends in Fig. 4, therefore, suggest that alloys having C+N contents

less than 0.128 wt.% behave similarly to AISI 304L, since their yield

stresses below 200 K are lower than expected. The same effect is

responsible for the break from the linear dependences at 195, 76, and

4 K in Fig. 3 for alloys having less than 0.128 wt.% carbon plus nitrogen.

Figure 5 shows that the ultimate tensile strengths for the steels

tested fall in the ranges 596 to 740 MPa at 295 K, and 1454 to 1735 MPa

at 4 K. The ultimate tensile strength increases slightly with increasing

interstitial content, but the effect is minor compared to the effects of

interstitials on yield stress, or compared to the effect of cryogenic

temperatures on yield and ultimate stresses. All of the tensile tests

at 4 K produced serrated stress-vs-strain curves, such as described by

Read and Reed^^^. This phenomenon Basinski originally attributed

to adiabatic heating. In the present study, the load drops vyere observed

to vary from 1.11 to 4.45 kN (250 to 1000 lbs), increasing with alloy

yield stress.

As shown in Fig. 6, the elongation and reduction of area measurements for all nine steels fall in the ranges 45 to 72% at 295 K and 30

to 48% at 4 K. These values are typical for commercial stainless steels

of similar composition. As listed in Table II, some of the steels of

24
this study show a progressive decrease of ductility at temperatures

between 295 and 4 K, whereas others exhibit an initial increase of

ductility between 295 and 200 K, followed by the decrease. In any case,

the level of ductility for the interstitial ly strengthened steels

remains useful at temperatures as low as 4 K.

Fatigue

Fatigue resistance at cryogenic temperatures decreases with increasing C+N content. As shown in Fig. 7, the resistance to crack initiation

at 77 K is reduced by about a factor of 4 for the conditions chosen in

these tests. The fatigue crack growth rates shown in Fig. 8 also

demonstrate that the da/dN values at 4 K and at a given aK value generally increase with increasing C+N content. The results agree with

existing trends for commercial AISI 304 L, AISI 304, and AISI 304 N

stainless steels, which show that the fatigue crack growth rates

increase with increasing nitrogen content.

Fracture Toughness

Jt results were obtained for seven of the austenitic stainless

Ic

steels tested, but the data for the two lower-strength heats with C+N

contents of 0.067 and 0.097 wt.% were invalid. The maximum J values in

these tests exceeded the allowable limits for 24.5-mm thick specimens,

the resistance curves deviated from the blunting line at exceptionally

low J values, and the curves could not be analyzed according to conventional techniques because there was no obvious intersection point at

the blunting line-resistance curve intersection. The data for the

seven other alloys tested at 4 K are listed in Table III, and the

corresponding estimates are plotted in Figs- 9 and 10.

25
The estimates from J-integral data were obtained from the

relationshi p

d-v')

where E is Young's modulus and v is Poisson's ratio.

The elastic constants for the stainless steels at 4 K were estimated at

E = 206.8 GPa and v = 0.30. The values of estimates ranged from 337

1 72

to 123 MPa-m for carbon plus nitrogen contents betv.'een 0.128 and

0.325 wt.%. As shown in Fig. 9, the estimates decreased linearly

with interstitial concentration. In Fig. 10, the estimates are also

shown to be inversely related to yield stress. The expression governing

the relationship between fracture toughness and yield stress is:

K. = 500 - 0.3

i c y

1 /2 This expression gives to within ±20 MPa-m '

. The fracture toughness results compare well with existing data, some of which are shown

for comparison in Fig. 10. Nitronic (Fe-Cr-Ni-Mn-N) alloys tend to have

higher nitrogen contents, which results in higher 4 K strength and lower

toughness compared with the AISI 304 alloys^^^. The AISI 304 N data of

[21

Read and Reed*- fall slightly below the present data trend, reflecting

differences owing to material variability or measurement uncertainty.

26
Martensitic Transformations

The approximate amount of a' (bcc) martensite formed during fatigue

and fracture at 4 K is plotted as a function of interstitial content in

Fig. 11. These results provide only an approximate indication of the

extent of phase transformation, because the measurements were taken

directly from the compact specimen fracture surfaces, which are rough and

unmachined. Some e (hep) martensite is known to form in addition to a'

phase, but this was not measured. Nevertheless, it is clear from Fig. 11

that the stability of the austenite is significantly increased by carbon

and nitrogen additions, but some martensite forms in all of the alloys

when fatigued or fractured at 4 K. The amount of martensite formed

increases with the plastic zone size, so more a' is detected on fracture

surfaces compared with fatigue surfaces, and more martensite forms in

the low-strength grades. The role of martensite transformation and its

effect on mechanical property results needs to be further investigated.

SUMMARY AND CONCLUSIONS

1 . Fatigue crack growth and fracture toughness measurements at 4 K

and conventional tensile properties at 295, 195, 76, and 4 K are

reported for Fe-18Cr-10Ni austenitic stainless steels having C+N

contents ranging from 0.067 to 0.325 wt%.

2. Carbon and nitrogen both increase the yield stress, but nitrogen

additions are much more effective. The yield stress for the alloys

tested increased from 222-358 MPa at room temperature to 329-1177 MPa

at 4 K. At room temperature the yteld stress tncreased linearly

with carbon plus nitrogen content. At low temperatures, the

27
yield stress also increased linearly for C+N contents gre^(ter than

0.128 wt.%

2. Fatigue crack growth resistance at 4 K decreased with increasing

C+N content.

4. Fracture toughness tests were performed using the single specimen

J-integral unloading compliance technique and 24.5-mm thick compact

specimens of the TL orientation. The tests yielded results

and Kj^ estimates for seven of the alloys at 4 K. The fracture

toughness decreased linearly with increasing C+N content, and the

Kj^ estimates at 4 K were inversely related to yield stress at 4 K,

with values ranging from 123 to 336 MPa-m^^^.

5. The mechanical properties of austenitic stainless steels vary

predictably with C+N content so that the 4 K properties can be

tailor-made by adjusting the interstitial element concentration.

ACKNOWLEDGMENTS

D. T. Read of NBS assembled the computer-aided, J-integral test

apparatus. D. T. Read, T. A. Whipple, A. C. Castille, D. S. Burkhalter,

and D. H. Beekman assisted in performing the tensile and fracture tests.

The stainless steel plates were supplied by the United States Steel

Corporation.

28
REFERENCES

1. R. P. Reed, F. R. Fickett, M. B. Kasen, and H. I. McHenry, "Magnetic

Fusion Energy Low Temperature Materials Program: A Survey," in:

Material s Studies for Magnetic Fusion Energy Appl i cations at Low

Temperatures-I , NBSIR 78-884, F. R. Fickett and R. P. Reed, Eds,

National Bureau of Standards, Boulder, CO, April 1978, pp. 245-335,

2. D, T. Read and R. P. Reed, "Fracture and Strength Properties of

Selected Austenitic Stainless Steels at Cryogenic Temperatures ," in:

Material s Studies for Magnetic Fusion Energy Appl i cations at Low

Temperatures-I

, NBSIR 79-1609, F. R. Fickett and R. P. Reed, Eds,

National Bureau of Standards, Boulder, CO, June 1979, pp. 81-122.

3. D. T. Read and R. P. Reed, "Toughness, Fatigue Crack Growth, and

Tensile Properties of Three Nitrogen-Strengthened Stainless Steels

at Cryogenic Temperatures , " in: Materials Studies for Magnetic

Fusion Energy Appl ications at Low Temperatures-I , NBSIR 78-884,

F. R. Fickett and R. P. Reed, Eds, National Bureau of Standards,

Boulder, CO, April 1978, pp. 91-154.

4. D. T. Read, R. P. Reed, and R. E. Schramm, "Low Temperature Deformation of Fe-18Cr-8Ni Steels," in: Materials Studies for Magnetic

Fusion Energy Appl ications at Low Temperatures-I

, NBSIR 79-1609,

F. R. Fickett and R. P. Reed, Eds, National Bureau of Standards,

Boulder, CO, June 1979, pp. 151-172.

5. D. T. Read and R. P. Reed, "Heating and Strain Rate Effects in

AISI 304L Stainless Steel at 4 K," in: Materials Studies for

Magnetic Fusion Energy Appl ications at Low Temperatures-I

, NBSIR 79-

1609, F. R. Fickett and R. P. Reed, Eds, National Bureau of Standards,

Boulder, CO, June 1979, pp. 125-147.

29
6. R. L. Tobler and R. P. Reed, "Tensile and Fracture Behavior of

Ni trogen-Strengthened , Chromium-Nickel -Manganese Stainless Steel at

Cryogenic Temperatures," Symposium on Elastic-Plastic Fracture,

American Society for Testing and Materials, Atlanta, GA, Nov. 1977.

7. P, P. Dessau, "LN^ and LH^ Fracture Toughness of Armco Alloy 22-13-5,

Report No. TID/SNA-2083 , Aerojet Nuclear Systems Co, Sacramento, CA

1971

8. D. C. Larbalestier and D. Evans, "High Strength Austenitic Stainless Steels for Cryogenic Use," Proceedings of Sixth International

Cryogenic Engineering Conference , Grenoble, France (1976), IPC

Science and Technology Press, Guildford, Surrey, England, pp. 345-347

9. R. B. Gunia and G. R. Woodrow, "Nitrogen Improves Engineering

Properties of Chromium-Nickel Stainless Steels," J . Mater. , JMLSA ,

5, June 1970, pp. 413-430.

10. R. Voyer and L. Weil, "Tensile and Creep Properties of a High

Nitrogen Content 18/10 (AISI 304L) Stainless Steel at Cryogenic

Temperatures," Adv. Cryo . Eng . 11 , 1965, pp. 447-452.

O 11. L.-A. Norstrbm, "The Influence of Nitrogen and Grain Size on Yield

Strength in Type AISI 316L Austenitic Stainless Steel," Met .

Sci . 11, 1977, pp. 208-21 2.

12. M. K. Booker and V. K. Sikka, "Effects of Composition Variables on

the Tensile Properties of Type 304 Stainless Steel," ORNL-5353,

Oak Ridge National Laboratory, Oak Ridge, TN, 1977, pp. 38.

13. P. Soo and W. H. Horton, "The Effect of Carbon and Nitrogen on the

Short-term Tensile Behavior of Types 304 and 316 Stainless Steels,"

Ifard NA 3045-2, Westinghouse Electric Corp., Madison, PA, pp. 48.

14. K, J. Irvine, D. T. Lewellyn, and F. B. Pickering, "High Strength

Austenitic Stainless Steels," J. Iron Steel Inst . 199 , Oct. 1961,

pp. 1 53-1 75.

30
15. K. J. Irvine, T. Gladman, and F. B. Pickering, "The Strength of

Austenitic Stainless Steels," J, Iron Steel Inst. , 207 , July 1969,

pp. 1017-1028.

16. C. E. Spaeder, Jr., and W. F. Domis, "Cryogenic and ElevatedTemperature Mechanical Properties of a High-Nitrogen Type 304

Stainless Steel," Report No. 44(01 2)-062-2 , United States Steel

Corporation, Pittsburgh, PA, 1969.

17. R. P. Reed, "A Cryostat for Tensile Tests in the Temperature

Range 300 to 4 K," Adv . Cryo . Eng

_1 1961 , pp. 448-454.

18. C. M. Fowlkes and R. L. Tobler, "Fracture Testing and Results for a

Ti-6A1-4V Alloy at Liquid Helium Temperature," Eng . Frac . Mech . 8

1976, pp. 487-500.

19. Standard Method of Test for Plane Strain Fracture Toughness of

Metallic Materials (Designation E399-74), 1974 Annual Book of ASTM

Standards , Part 10, 1974, pp. 432-451.

20. G. A. Clarke, W. R. Andrews, P. C. Paris, and D. W. Schmidt,

"Single Specimen Tests for Jj^ Determination," in: Mechanics of

Crack Growth , ASTM STP 590, American Society for Testing and

Materials, Philadelphia, PA, 1976, pp. 27-42.

21. J. G. Merkle and H. T. Corten, "A J-Integral Analysis for the

Compact Specimen, Considering Axial Force as Well as Bending

Effects," Pressure Vessel Tech ., Trans ASME, 1974, pp. 1-7.

22. J. D. Landes and J. A. Begley, "Test Results from J-Integral

Studies: An Attempt to Establish a Jj^ Testing Procedure," in:

Fracture Analysis, ASTM STP 560, American Society for Testing and

Materials, Philadelphia, PA, 1974, pp. 170-186.

31
23 . R. P. Reed and R. P. Hi Resell, "The Stability of Austenitic Stainless Steels at Low Temperatures as Determined by Magnetic Measurements," Adv. Cryo. Eng. 4 1960, pp. 84-100.

24. R. P. Reed and C. J. Guntner, "Stress-Induced Martensitic Transformations in 18Cr-Ni Steel," Trans. Met. Soc. AIME 230, 1964,

pp. 1713-1720.

25. R. L. Tobler, R. P. Reed, and D. S. Burkhalter, "Temperature

Dependence of Yielding in Austenitic Stainless Steels," in this

volume.

26. Z. S. Basinski, "The Instability of Plastic Flow of Metals at Very

Low Temperatures," Proc. Roy. Soc. A240 , 1957, pp. 229-242.

27. Z. S. Basinski, The Instability of Plastic Flow of Metals at Very

Low Temperatures," Aust . ^J. Phys . 1

, I960, pp, 354-353.

32
LIST OF TABLES

Table No.

I. Stainless steel compositions (wt.%), hardnesses, and grain sizes.

II. Tensile property results (one test per temperature)

III. J-integral toughness results at 4 K.

33
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Tensile

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(one

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per

temperature)

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35
Table III. Yield stress and J-Integral toughness results at 4 K.

Specimen C+N Content Yield Stress, ay

Flow Stress, 25 J.

Ic 'hIc

1/ Ic Estimates*

(wt.%) (MPa) (MPa)

of

(mm) (kJ/m^) (MPa-m^/^)

1 0.067 329 892 NA NA NA

2 0.097 445 996 NA NA NA

3 0.28 530 1029 12.2 499 337

4 0.157 745 1158 6.6 312 266

5 0.187 876 1288 4.6 230 230

6-A 0.214 896 1264 4.3 218 222

6-B 0.214 896 1264 5.3 272 249

7 0.270 1186 1411 11.9 67 124

8-A 0.297 1178 1456 1.5 88 141

8-B 0.297 1178 1456 1 .3 78 133

9 0.325 1286 1496 1.1 63 124

2 2 *Calculated from the expression • E/(l-v ), where it is assumed that

E - 206.8 GPa and v = 0.30 at T = 4 K.

36
LIST OF FIGURES

Figure No.

1. Specimen geometries used for conventional tensile and fracture

mechanics tests.

2. Typical J-resistance curve for an austenitic stainless steel tested

at 4 K.

3. Dependence of yield stress on carbon and nitrogen level, with

temperature as a parameter.

4. Temperature dependence of yield stress. Numbers refer to heat numbers. Table I.

5. Dependence of ultimate tensile stress on carbon and nitrogen, with

temperature as a parameter.

6. Temperature dependence of elongation and reduction of area.

7. Fatigue crack initiation resistance at 77 K.

8. Dependence of fatigue crack growth rates at 4 K on carbon and

nitrogen content.

9. Dependence of at 4 K on carbon and nitrogen content.

10. Inverse relationship between fracture toughness and yield stress

for austenitic stainless steels at 4 K.

11. Relative amounts of a' (bcc) martensite formed during tests of

compact specimens at 4 K.

37
COMPACT SPECIMEN

Thickness = 24.4 mm

Thickness = 2.54 mm

Figure 1. Specimen geometries used for conventional tensile and

fracture mechanics tests.

33
1200

1000

J-Resistance Curve

Fe-l8Cr-l0Ni Stainless Steel

(C+N = 0.128 weight percent)

CNJ

<D

O

o>

(D

800

600

400

200

0

Compact Specimen

25.4 mm thick

a/W« 0.630

TL Orientation

Blunting Line

J = 2AacTf

-I 0 12 3

Crack Extension, Aa (mm)

Figure 2. Typical J-resistance curve for an austenitic stainless

Steel tested at 4 K.

39
Figure 3. Dependence of yield stress on carbon and nitrogen level,

with temperature as a parameter.

40

Yield

Stress,

ay

(ksi)
Yield

Stress,

<Tw

(MPa)

2000

Fe-l8Cr-IONi

100 ^0

I I 100 200 300

Temperature (K)

Figure 4. Temperature dependence of yield stress. Numbers refer to heat

numbers. Table I.

41
Ultimate

Tensile

Strength,

<Tw

(MPa)

Carbon + Nitrogen (weight percent)

Figure 5. Dependence of ultimate tensile stress on carbon and nitrogen,

with temperature as a parameter.

42
90

80

70

60

50

40

30

20

Temperature (K)

gure 6. Temperature dependence of elongation and reduction of area.

43
Fe-l8Cr-IONi

CT

Specimens

0 0 0 0 0 0

0 0 0 0 0

10 ro OJ —

ijODJO UJUU 0'| D 84D!;iU| 04 S8|0A00|!>j

44

0.10

0.20

0.30

Carbon

+

Nitrogen

(weight

percent)
Fatigue

Crack

Growth

Rate,

da/d N

(mm/cycle)

Figure 8. Dependence of fatigue crack growth rates at 4 K on carbon and

nitrogen content.

45
400

CVJ

O 300 0.

"D

I—

</> CO

a>c

3

0>

w.3

O

O

200

100

0.10 0.20 0.30

Carbon + Nitrogen (weight percent)

Figure 9. Dependence of at 4 K on carbon and nitrogen content

46
Fracture

Toughness,

Kj(j(J)

(MPaYield Stress, Oy (MPa)

Figure 10. Inverse relationship between fracture toughness and yield stress

for austenitic stainless steels at 4 K.

47
Relative

Amount

of

BCC

Martensite

Figure 11. Relative amounts of a' (bcc) martensite formed during tests

of compact specimens at 4 K.

48
TEMPERATURE DEPENDENCE OF YIELDING IN

AUSTENITIC STAINLESS STEELS

National Bureau of Standards

49

TEMPERATURE DEPENDENCE OF YIELDING IN

AUSTENITIC STAINLESS STEELS*

R. L. Tobler, R. P. Reed, and D. S. Burkhalter

National Bureau of Standards

Boulder, Colorado

ABSTRACT

The temperature dependence of yielding is discussed with specific

reference to austenitic stainless steels. Tensile yield stress (o ) measurements at 0.2% plastic strain, a strain rate of 0.02 mm/s, a)id

temperatures between 4 and 295 K are reported for a heat of AISI 316

stainless steel. The results are compared with previously published

data for the same heat at elevated temperatures , and the yield stress

is observed to decrease from 660 MPa at 4 K to 102 MPa at 922 K. The

decrease is regular, with no obvious anomalies, and apparently no low

temperature martensitic transformations at 0.2% strain. The merits

and limitations of various expressions for describing the temperature dependence of the yield stress are then evaluated, and an excellent data

fit is obtained using the expression;

^®o“^A^ exp(-CyT)

where is an athermal stress, a is the yield stress at absolute zero,

and Cy IS a constant related to t’^ermodynamic parameters for thermally

activated plastic flow.

Key Words: Cryogenic properties of materials; iron alloys; low temperatu

mechanical properties; plastic deformation; temperature

effects; stainless steel alloys; yield stress.

*Supported by the Office of Fusion Energy, U.S. Department of Energy.
INTRODUCTION

The tensile properties of stable austenitic stainless steels, such

as AISI 310 ( Fe-25Cr-20Ni-2Mn-0. 08C) , have a regular temperature dependence with trends similar to the classical behavior expected for metals

and alloys having the face-centered cubic structure [1]. The classical

behavior consists of a smooth, monotonic increase of yield stress as

temperature is lowered to about 60 K. In contrast, metastable stainless

steels, such as AISI 304L (Fe-1 8Cr-9Ni-2Mn-0.03C) , have yiel d-stress-(a^)-

vs. -temperature anomalies that are believed to be associated with low

temperature martensitic transformations to body-centered cubic (a') and

hexagonal close-packed (e) phases [2], It has been suggested that such

transformations may lower the 0.2% yield stress [2]. In addition, both

the stable and metastable alloys are capable of anomalous behavior at

extreme cryogenic temperatures (below 60 K) owing to deformation mode

changes associated with dislocation tunneling or magnetic transitions [3].

The metastable stainless steel AISI 316 ( Fe-1 8Cr-l 3Ni-2Mn-2Mo0.06C) remains austenitic after cooling to cryogenic temperatures but

transforms to approximately 50% a' martensite when strained to fracture

at 4 K [4]. Since AISI 316 is an important candidate for superconducting

machinery and magnetic fusion, its temperature dependence of yielding

was investigated. As described below, anomalous effects on the a^-vs.-

temperature relationship owing to the possible formation of martensite

at 0.2% strain proved negligible in our tests. Therefore the data are

used to appraise various analytical expressions for describing the

temperature dependence of for AISI 316 and related alloys including

AISI 310.

52
EXPERIMENTAL PROCEDURE

Oak Ridge National Laboratory donated the AISI 316 stainless steel

that Sikka F61 had tested previously at elevated temperatures. This

material was a 25.4-mm thick plate with the following composition in

wt.%: Fe-17.25Cr-13.48Ni-l .86Mn-2.84Mo-0.02P-0.019S-0.58Si-0.057C-0.030N.

Twenty flat, pin-loaded tensile test specimens in the transverse

orientation were machined from stock. The specimens had a 2.5-mm by

4-mm rectangular reduced cross section and a 25-mm gauge length. After

machining, the specimens were annealed at 1289 K in a vacuum tube furnace.

After 1 h, the tube was removed from the furnace, the vacuum was broken,

and the specimens were cooled in air. Following this treatment, the

average grain diameter was 67 ym (ASTM grain size No. 4.9), and the

hardness was Rockwell B 75.

Tensile tests used the screw-driven machine and the cryostat

described by Read, Reed, and Schramm [2]. The strain rate was 0.02 mm/s.

The specimens were instrumented with a double-beam strain gauge extensometer, as well as platinum or carbon resistance thermometers. Tests at

4 and 77 K were performed by direct immersion in liquid helium or

liquid nitrogen. Other temperatures were achieved by controlled thermal

conduction using liquid helium or nitrogen reservoirs and heating wires

attached to the specimen grips and load train. Current to the heating

wires was varied until thermal equilibrium was reached at the desired

test temperature. Load-vs. -deflection curves were recorded using commercial load cell and extensometer outputs, and the yield stress at 0.2%

plastic strain was calculated in the conventional manner. Two or three

specimens were tested at 295, 76, and 4 K. Individual measurements did

53
not deviate more than ±6% from the average values at these temperatures.

The specimens were plastically strained only O.Z%; they were then unloaded

and examined using a torsion bar magnet device [2] to check for the

presence of ferromagneti sm and, hence, a' martensite at this small

deformation.

RESULTS AND DISCUSSION

Seventeen yield stress measurements were obtained at 13 temperatures between 4 and 295 K. These measurements are listed in Table I,

along with Sikka's elevated temperature data for the same heat. The

combined results represent the most extensive yield stress data available for a single heat of austenitic stainless steel. The two data sets

are in good agreement at room temperature where they overlap: Sikka's

yield stress of 216 MPa at 298 K is only S% higher than the average

value of 205 MPa at 295 K from this study. The yield stress vs temperature plot for this heat of AISI 316 (shown in Fig. 1) indicates a

regular monotonic decrease of from 660 to 102 MPa as temperature is

raised from 4 to 922 K. The decrease is most dramatic between 4 and

295 K, amounting to a factor of 3.2. Selected specimens from this

temperature range were examined at room temperature using torsion bar

magnet and metal 1 ographi c techniques, but measurable amounts (0.1% or

greater) of the ferromagnetic a' phase could not be detected after

plastic strains of 0.2%. This search for martensite was cursory, and an

x-ray inspection would be necessary to confirm whether small amounts of

a' or e phases had formed; but the absence of any obvious yield stress

anomalies implied that the martensitic transformation had a negligible

effect on in this heat of AISI 316 stainless steel.

54
Sanderson and Llewellyn's data for AISI 310 and AISI 316 stainless

steels [6] are also plotted for comparison in Fig. 1. The trend for

AISI 310 (77 5 T 5 800 K) resembles the trend for the AISI 316 alloy of

this study. This is significant because AISI 310 is highly alloyed and

completely stable with respect to low temperature martensitic transformations. Thus, neither in the AISI 310 of Sanderson and Llewellyn's

study nor in the AISI 316 of this study are there any indications of

anomalous yield stress decreases at temperatures between 200 and 100 K,

such as those observed in AISI 304L where strain-induced martensitic

transformations occur.

On the other hand, the Sanderson and Llewellyn data for AISI 316

begin to deviate from the present results at temperatures below 176 K,

and the difference increases at still lower temperatures amounting to

100 MPa at 77 K. Sanderson and Llewellyn's heat (Fe-16.72Cr-ll.27Ni1 .52Mn-2.58Mo-0. 32Si-0.05C) had a lower alloy content than the one o.

this study (Fe-17.25Cr-13.48Ni-l.86Mn-2.34Mo-0.58Si-0.057C-0.03N), which

may explain the lower yield stress, since low-temperature martensitic

transformations are believed to lower the yield stress, and such transformations would be more probable for their heat. This suggests the

possibility that measurements at strains higher than 0.2%, where martensite products are formed, might reveal flow-stress-vs. -temperature

anomalies in the heat of AISI 316 of this study. This possibility will

be investigated in future tests.

CURVE-FITTING TECHNIQUES

Various expressions have been proposed to specify the temperature

dependence of yielding for alloys exhibiting regular behavior. To

55
evaluate their merits and limitations, data spanning a broad temperature

range are needed. Table I fills this need for AISI 316 stainless steel.

In the following analysis, yield stress, a^, is in MPa and temperature

T, is in kelvin.

Polynomial

The most common engineering expressions for general curve-fitting

purposes are polynomials of the form:

Oj, = Cjj + C^T + C^-\^ + + .... (1)

The number of terms used may vary, but by mathematical regression it is

possible to determine the coefficients (C^, C-j , C

2

» •••)» thereby fitting

practically any data, regardless of alloy or temperature range. For

AISI 316, the following four-term solution was obtained:

= 680.5 - 2.468T + 3.534 x 10"^T^ - 1.670 x lO'^T^ (2)

which predicts to within 8% of the measured values for the temperatures listed in Table I. Polynomials are a useful design tool because

of their wide applicability. To their disadvantage, however, polynomials contain terms that have no physical significance or theoretical

basis, and extrapolations of polynomials beyond the actual data accumulation range are not accurate.

Yang's Expression

Based on empirical data and some theoretical support, Yang [7]

proposed the expression:

Oy = CjT ® (3)

56
where and Cg are constants for a given alloy. If this expression is

valid, data in agreement with it should exhibit a linear trend when

using this equation to describe his yield stress measurements for highstrength iron, nickel, aluminum, and titanium alloys at temperatures

between 20 and 300 K. But the data of annealed AISI 316 do not fit

Eg. (3) (Fig. 2). Poor fits are also found using annealed AISI 310

data [6,8], thus, Eq. (3) is Inapplicable to annealed austenitic stainless

steel data.

Zener and Hollomon's Expression

Building on Becker's thermally activated plastic-flow theory

[9],- Zener and Hollomon [10] proposed the following expression for a

pearl i tic steel between 293 and 77 K:

where a and r are constants, Q is a heat of activation, and R is the

universal gas constant (R = 8.31 J/mol K) . According to this expression.

steel [10], molybdenum [11], and tungsten [12]. As indicated in Fig. 3,

the Zener and Hollomon expression also holds for AISI 316 and AISI 310

stainless steels, but only at elevated temperatures . The governing

equation for AISI 316 between 300 and 922 K is:

plotted as log vs. log T. Yang claimed satisfactory results when

= a exp(rQ/RT) (4)

Oy ^ 72.9 exp(2560/RT) (5)

57
Elevated temperature data for AISI 304 and 310 stainless steels [6,13]

fit similar expressions, but the r and Q parameters cannot be evaluated

separately unless variable strain rate tests are performed. Another

obvious disadvantage of the Zener and Hollomon expression is that it

fails for austenitic stainless steels at cryogenic temperatures , as

shown clearly in Fig. 1

Yaroshevich and Ryvkina's Expression

A very successful expression stems from classical theory which

relates the yield stress to two terms [14]:

Here is an athermal (temperature independent) stress determined by

metallurgical structure, and is a thermal stress controlled by lattice

vibrations and thermal ly-acti vated rate processes. From thermodynamic

and thermally activated plastic flow theories, Yaroshevich and Ryvkina [15]

developed the following expression for a.^:

In this expression, is the extrapolated yield stress at absolute

zero, is the value of a-j- at absolute zero, and Cj is a physically

significant material constant related to activation energy, entropy, and

strain rate [15]. A complete expression for the yield stress then

fol 1 ows

( 6 )

Of = {^Q-Oi:^)exp{-CyT) (7)

" ^A

^ expl-C^T) ( 8 )

58
or:

log(a^-a^) = log(a^-a^)-C

7 T (9)

where a„ is a constant that can be estimated from the a -vs-temperature

^ y

plot for a given alloy.

If Eq. (9) holds, a plot of log(a -o.) vs. T should be linear, and

y ^

it is for AISI 316 in Fig. 4.. From such a plot, is the ordinate

intercept at T = 0, and Cy is the slope. From Fig. 1, Op^ for AISI 316

is approximately 100 MPa. Thus, the following solution is obtained for

AISI 316:

= 100 + 676 exp(-6.16xlO"^T) (10)

Between 922 and 30 K, Eq. (10) predicts the data of Table I to

within 20 MPa. Below 30 K, the fit is less accurate, and the measured

values fall below Eq. (10) predictions by as much as 115 MPa. This

deviation at extreme cryogenic temperatures may be a subtle manifestation of paramagnetic-to-antiferromagnetic transformations and dislocation tunneling effects, which represent departures from the assumptions

underlying Eqq. (8), (9), and (10). Magnetic transition (near 60 K and

below) and dislocation tunneling (near 4 K) are both expected to reduce

the yield stress, according to Skibina's [3] work on similar steels.

Moreover, Ledbetter's elastic property measurements for the AISI 316 of

this study confirm that Young's modulus is anomalous below 60 K, and

therefore, the magnetic transition probably occurs in this alloy [16].

59
The Yaroshevich and Ryvkina expression also works well for AISI 310

at temperatures between 77 and 800 K. In terms of Eq. (8), the strength

difference between AISI 316 and 310, as illustrated in Fig. 1, may be

ascribed primarily to the difference in From Fig. 1 the values

for AISI 310 and 316 are 183 and 100 MPa, respectively, whereas

Fig. 4 shows that the terms are nearly identical. This result

implies that the thermally activated flow processes for these two

congeneric alloys are similar.

Although the Yaroshevich and Ryvkina expression is successful for

austenitic stainless steels, some practical difficulties remain. A

large data base, including elevated temperatures , is necessary to determine parameters such as accurately. Rarely are elevated temperature

data available for cryogenic alloys, and rarely are cryogenic data

available for elevated temperature alloys.

A Simplified Cryogenic Expression

It was discovered that the AISI 316 data of this study between 4

and 295 K have a linear relationship when log is plotted vs. T.

Using this format, the data for AISI 316 and 310 are compared in Fig. 5.

Below 295 and 250 K, respectively, the data for both alloys obey an

equation of the form:

Oy = Oq exp(-CgT) (11)

Here is the extrapolated yield stress at absolute zero, and Cg is an

empirical constant without theoretical significance corresponding to the

slope of the straight line in Fig. 5. Solving these parameters for

AISI 316 gives:

60
( 12 ) = 697 exp(-4. 08x10 ^T)

The fit between 4 and 295 K is such that none of the experimental data

vary by more than 1 % from the Eq. (12) predictions, and most data agree

to within 4%, which is the order of specimen-to-specimen variability.

A practical use of Eq. (11) is illustrated in Fig. 6, where selected

data for other austenitic stainless steels also fit the proposed format

of Eq. (11). The data shown for AISI 310 [8,17,18], Fe-21 Cr-6Ni -9Mn [19],

Fe-18Cr-3Ni-13Mn [20], and alloy A-286[21] suggest the possibility of

predicting the 4 K yield stress from only 77 and 295 K data. This is

reasonable for stainless steels that are known to remain austenitic

after 0.2% plastic strain at cryogenic temperatures , but broader applicability is questionable. For example, other yield stress data

[6] for AISI 316 at 0.2% plastic strain are not linear when plotted as

log vs. T, and Eq. (11) should not be expected to hold if a change in

deformation mechanism occurs at cryogenic temperatures

The latter point is emphasized in Fig. 7, where values at 0.2%

plastic strain for AISI 310 and 304L are compared on a log vs. T

scale. Warren and Reed's data [8] for the stable AISI 310 indicate a

trend similar to the present results for AISI 316, closely fitting

Eq. (11). But AISI 304L has a lower austenite stability compared with

AISI 316 or 310, so spontaneous or stress-induced martensitic phase

transformations are expected for AISI 304L. The data of Read et al

[2] for AISI 304L show the typical strengthening between 300 and 200 K,

but anomalous behavior is observed between 200 and 100 K where ay

decreases

61
Below 100 K, increases again, following an exponential function

similar to Eq. (11). This anomaly could be directly associated with a

partial phase transformation to e martensite, as remarked elsewhere in

this paper.

Figures 5 to 7 indicate that Eq. (11) is a good approximation of

stable austenitic stainless-steel data in the cryogenic range, opening

prospects of a parametric analysis whereby mechanical and metallurgical

effects might be characterized in terms of the parameters and Cg.

For example, the effects of cold work on the behavior of AISI 310

[18] are illustrated in Fig. 8 As cold work increases, increases

and Cg decreases. Eq. (8) could also be used for parametric analysis,

but Eq. (11) is simpler. The disadvantages of Eq. (11) compared with

Eq. (8) are that Eq. (11) is restricted to low temperatures and the

parameter Cg lacks the physical significance attached to Cy,

SUMMARY AND CONCLUSIONS

The yield stress was measured for one heat of AISI 316 stainless

steel at 0.2% plastic strain and at temperatures between 4 and 295 K.

Comparison with el evated-temperature data showed that the yield stress

decreased rapidly from 660 MPa at 4 K to 102 MPa at 922 K. The decrease

was regular and monotonic, with no obvious anomaly, and no measurable

amount of a' martensite was detected after 0.2% plastic strain. But it

is suspected that anomalies in the yield or flow stress temperature

dependences may occur in heats of AISI 316 at lower alloy contents or at

higher strains.

Since the yield-stress-temperature relationship for the AISI 316

steel of this study was regular, a number of curve-fitting techniques

62
were evaluated. Various expressions describing the temperature dependence of for AISI 316 and other stainless steels that remain austenitic at 0.2% strain are summarized in Table II. The most satisfactory

expression (Eg. 7) for the 922-degree-temperature-range considered is

derived from thermodynamic and thermally activated olastic-flow theories

Oy = 0 ^

+ exp(-CyT),- where 0^ and Cy are material constants and

0

^ is the extrapolated yield stress at absolute zero. For AISI 316,

the experimental agreement with this expression is excellent, except for

subtle deviations below 30 K that may be attributable to magnetic

transition and dislocation tunnelling effects that moderately reduce

below expected values.

ACKNOWLEDGMENT

D. Baxter of the National Bureau of Standards, Boulder, Colorado

contributed by performing tensile tests and data reduction. Her assistance is gratefully acknowledged.

63
NOTATION

Roman Symbols

^o’ ^1’ ^2

~ coefficients or exponents used in expressions for the temperature dependence of yield stress.

Q

= heat of activation.

r = constant used in Eq. 4. R = universal qas constant.

T = temperature in degrees Kelvin.

Greek Symbols

o = constant used in Eq. 4, having units of stress.

= the athermal component of yield stress.

= the yield stress at absolute zero, obtained by extrapolation.

Oj = the thermal component of yield stress.

= the yield stress at 0.2% plastic strain.

64
REFERENCES

1. K. G. Brickner and J. D. Defilippi, in Handbook of Stainless Steels ,

(D. Peckner and I. M. Bernstein, eds.)> McGraw-Hill Book Company,

New York (1977), p. 20-1.

2. D. T Read, R. P. Reed, and R. E. Schramm, in Materials Studies for

Magnetic Fusion Energy Applications at Low Temperatures - II,

NBSIR 79-1609, National Bureau of Standards, Boulder, Colorado

(June 1 979) , p. 1 51 .

3. L. V. Skibina, Prob. Prochn. 2:71 (1976).

4. D. C. Larbalestier and H. W. King, Cryogenics 13:160 (1973).

5. V. K. Sikka, "Product Form Characterization of Reference Heat of

Type 316 Stainless Steel," Report No. ORNL-5384, Oak Ridge National

Laboratory, Oak Ridge, Tennessee, (1978).

6. G. P. Sanderson and D. T. Llewellyn, J. Iron Steel Inst. 207:1129

(1969).

7. C. T. Yang, J. Eng. Ind. 88:117 (1966).

8. K. A. Warren and R. P. Reed, "Tensile and Impact Properties of

Selected Materials from 20 to 300°K," NBS Monograph 63, National

Bureau of Standards, Boulder, Colorado, (1963).

9. R. Becker, Z. Tech. Phys . 7:547 (1926).

10. C. Zener and J. H. Hollomon, J. Appl . Phys. 15:22 (1944).

11. J. H. Bechtold, J. Mets ., 5:1469 (1953).

12. J. H. Bechtold and P. G. Shewmon, Trans. Am. Soc. Met. 46:397 (1954).

13. M. K. Booker and V. K. Sikka, "Effects of Composition Variables on

the Tensile Properties of Type 304 Stainless Steel," ORNL-5353, Oak

Ridge National Laboratory, Oak Ridge, Tennessee (1977).

14. H. Conrad, in High Strength Materials (V. F. Zackey, ed.) John

Wiley & Sons, Inc., New York (1964), p. 436.

15. V. D. Yaroshevich and D. G. Ryvkina (1970). Sov. Phys. -Sol id

State (Engl. Transl .

12(2) (1970).

65
16 . H. M. Ledbetter, National Bureau of Standards, Boulder, Colorado,

unpublished data.

17. J. M. Wells, R. Kossowsky, W. A. Logsdon, and M. R. Daniel, in

"Materials Research for Superconducting Machinery - III," NBS/ARPA,

Semi-Annual Technical Report ADAOl 2365 (1975), p. 1.

18. J. L. Christian, J. D. Gruner, and L. D. Girton, Trans. Am. Soc. Met.

57:199 (1964).

19. R. L. Tobler and R. P. Reed, in Elastic-Plastic Fracture , ASTM STP 668,

American Society for Testing and Materials, Philadelphia (1979),

p. 537.

20. D. T. Read and R. P. Reed, in Materials Studies for Magnetic Fusion

Energy Applications at Low Temperatures - I , NBSIR 78-884, National

Bureau of Standards, Boulder, Colorado (April 1978), p. 91.

21. R. P. Reed, R. L. Tobler, and R. P. Mikesell, in Advances in Cryogeni c

Engineering. Vol . 22 , Plenum Press, New York (1977), p. 68.

66
LIST OF TABLES

Table 1. Yield stress versus temperature data for annealed AISI 316

plate.

Table 2. Summary of expressions for the temperature dependence of

yield stress, based on attempted data fits for AISI 316 and

related stainless steels.

67
Table I. Yield Stress vs. Temperature Data for Annealed AISI 316

Stainless Steel

Temperature, K 0.2% Yield stress,*MPa Temperature, K

0.2% Yield stress,^

MPa

4 663 298 216

4 656 366 174

30 599 477 135

54 578 589 119

76 512 700 115

76 516 811 103

95 493 922 102

123 440

153 390

193 317

214 299

218 292

242 266

250 248

295 213

295 193

295 209

*Data from this study.

^Data from Sikka [5].

68
Table

II.

Summary

of

Expressions

for

the

Temperature

Dependence

of

Yield

Stress,

Based

on

Attempted

Data

Fits

for

AISI

316

and

Related

Stainless

Steels. 1 i. OO 4- c cu 1 G 1 1— O -X •t— CD. CU +j ... E E u E (— E OO E •1- 03

1/) fO cu CU cu o cu 4-> 03 4-

CO 1—u +J .E (DO > 03 +-> O) U 1 o cn E 03 E oO +J 1— 00 •r- .C -»-> E 03 x: G • r— •o E 00 -M CD3 O E CD. G

S- S- c •t— •!— cu E 00 E 1— 4-J nD O) o c .c 1— .E CU +-> G X E

cn -i-> •1— cu oi +-> E E 4-> . 03

X O) +-> +-> T3 03 E G E X

S- O) fn 00

O

U 1— +J • 1— T3

G C/0 »- sz t— =3 03 • (— E 03 U G X G

•' 4-> o 03 1 (

M- E • r— 4-> 4-> E cn

cn Q. s CL CU (U 4- 03

O E

c: -M fD s- o CL in CD. 4- > G 4- 03 r- 13 U o >— 03 =3 E G G E E

S- JO 4-> C|- CU

O 1

•r— E

OJX CU 4-> TD +->

UG 4-

O

G

QJ ^ CU OO JC E E G •r— -r— C CU 1— 4-> JD 03 "O "O +-> T3 00 E

•t- cn-o •r— CU E E CO

G

cn c: c: 03 .d r. . >o+-> 03X

O

G cn

sc fO fC M- +-> >> E 03 +-> E o . cu s- O O > CO .. CL >> G

cu JD 1— E CU CU +-)

G

X E E c cu CJ CU 03 O .E 1— E U E G G E

•r- s.. c 1—

C

U 4-> (U 03 E 03

X

Z5 rO JD •1— 00 03 OO cn

G 03

CU -M

u 03 +-> CU E T3 00U CO E X E

</) 1— fO •r- C_) 00 cu E •>- E E *1

—

G •1-

X

G

JU S- M- •r— 1 E E 03

O 4-

u E CD. c 03 CU •r— i— CU O -E T3 u •I— G G E E

O)

O CD. c CD. CU CU 03 (U

U E E G >-G E E cn Q.4-> 00 o: E I/O -t— 1— cn E X E 1— CU •1— 03 00 cu *-> CU 03 E 03 •r" G •1— 00 o SO. +-> CO E cn CD. JD CU •I— 00 X cni— E u so. • •1— 00 E T3 E cn E E G O

03 S- f— O) 00 03 CU CU >)

OG 1 >0G

G

O

>)

O 03 4-5 >, cu E -M -M r- >,E> 03 03

X E

S- >1

U rO 1— 1

E OO E 03

G E G OO

fO •— >1 • 1—U 1— c CU

O "O E U E •t— 1

G E CU o CO 13 03 -1— E CD. CU O CO 03 CD. OO > E -O r- cj E 03 E CL -M E (U G >o+-> E 00 O *1— (

U OO +-> -M E 03 +-> x: X SZ (O •r- G X LT) IS CD Q- 03 =) 00 03 (O > CO +J 4-> CD.T3 C/3 1— 03 +

t—

1— 1

>> C_)

C_J 1 +

CO CD.

C hX o c^ cu

*r— c_> 00 + .—

-—

^—. LO CXI h— 1

O) 1— Cd

G CO

s_ CXI 1 CG

a. o Cr o 1 X + E G

UJ 1—

— ^ 

CD. 1

CO Cu

X o o X + G + 1— CU o CO car

O

C_) C_)

DG

D

II II II II II >> >0 G D D G D

S-<u -QE ^— c 4-> ^ . ^ . ^

(

X 1

CO CO 1— c cu - . '—

V >

o 4-> •r" . . • • 4-> c cr cr cr CD" cr

fO • 1— LjJ LjU LlJ LiJ LlJ 3

cr

UJ

69
LIST OF FIOURES

Figure 1

Figure 2.

Figure 3.

Figure 4.

Figure 5.

Figure 6.

Figure 7.

Figure 8.

Effect of temoerature on the yield stress of annealed AISI 316

plate, with literature data for comoarison.

Plot of log Oy vs. log T for AISI 316, illustrating the unsatisfactory fit to Eg. (3).

Plot of log Oy vs. 1/T, illustrating the aopl icabil ity of

Eq. (4) for AISI 316 and 310 alloys at elevated temoeratures

Plot of log (a^-a^) vs. T, illustrating the aool icabi 1 ity of

Eq. (7) for AISI 316 and 310 stainless steels.

Plot of log Oy vs. T, illustrating the applicability of Eq. (11)

for 316 and 310 stainless steels at cryogenic temperatures.

Plot of log Oy vs. T, illustrating the applicability of Eq.

(11) to data for various austenitic stainless steels at cryogenic

temperatures

Plot of log vs. T, illustrating the anomalous behavior of

AISI 304L compared with the more heavily alloyed and more

stable AISI 316 and 310.

Effect of cold work on the yielding of AISI 310, using the

format of Eq. (1 1)

70
YIELD

STRESS,

a

.

MPa

100

80

40

20

Figure 1. Effect of temperature on the yield stress of annealed AISI 316

plate, with literature data for comparison.

71

YIELD

STRESS,

a

,

ksi
Figure 2

YIELD STRESS, a , MPa

y

YIELD STRESS, a , ksi

y

. Plot of log a vs. log T for AISI 316, illustrating the un- satisfactory tit to Eq. (3).

72

1000
TEMPERATURE, K

1000 100

100

80

60

40

30

20

Figure 3. Plot of log a vs. 1/T, illustrating the applicability of

Eg. (4) for AiSI 316 and 310 alloys at elevated temperatures.

73

YIELD

STRESS,

a

,

ksi
YIELD

STRESS

-a.

,

a,

-100,

MPa

74
700

L S >1 ‘SS3yiS Q13IA

75

for

316

and

S'^O

stainless

steels

at

cryogenic

temperatures.
200

to

100

70

50

Figure 6. Plot of log a vs. T, illustrating the applicability of Eq.

(11) to data for various austenitic stainless steels at cryogenic

temperatures.

76

YIELD

STRESS
YIELD

STRESS,

a

,

MPa

800

0 100 200 300

TEMPERATURE, K

Figure 7. Plot of log ^

vs. T, illustrating the anomalous

AISI 304L compared with the more heavily alloyed

stable AISI 316 and 310.

behavior of

and more

YIELD

STRESS,

a

,

ksi
YIELD

STRESS,

a.

,

MPa

200

100

70

50

TEMPERATURE, K

Figure 8. Effect of cold work on the yielding of AISI 310, using the

format of Eq. (11 )

78

YIELD

STRESS,

a

,

ksi
ANOriALOUS LOW-TEMPERATURE ELASTIC-CONSTANT BEHAVIOR

IN Fe-13Cr-19Mn

National Bureau of Standards

79

ANOMALOUS LOW-TEMPERATURE ELASTIC-CONSTANT BEHAVIOR IN Fe-1 3Cr-l 9Mn*

H. M. Ledbetter, Research Metallurgist

Fracture and Deformation Division, National Bureau of Standards,

Boulder, Colorado 80303

Elastic constants determined by sound-velocity measurements in an

Fe-13Cr-19Mn steel show anomalous temperature behavior around 274 K,

indicating a paramagnetic-anti ferromagnetic transition.

Key Words: Elastic constants; iron alloy; Neel transition; stainless steel.

81
1-3

Recent studies show that austenitic stainless steels with Mn and

N substituted for Ni show anomalous physi cal -property/ temperature behavior

Ledbetter^ observed that Fe-21 Cr-6Ni-9Mn exhibits elastic-constant

anomalies when cooled to about 130 K, and conjectured that a paramagneticanti ferromagnetic (N6el ) transition caused the anomalies. Collings

2

and Ledbetter verified the N6el transition explanation in studying the

elastic constants and magnetic susceptibility of a similar alloy, Fe1 8Cr-3Ni-l 3Mn , which behaves anomalously upon cooling to about 190 K.

3

Neel -transition-induced anomalies occur also in Fe-22Cr-l 3Ni-5Mn at

temperatures near 33 K.

3 Primarily, Mn content controls the anomalous behavior . Higher Mn

means a higher Neel transition temperature, Tj^, and larger elasticconstant changes at

Studying the limiting case -- Mn and N substituted completely for

Ni -- becomes of considerable interest. Such an alloy should show a

high T|^ and a large, anomalous elastic-constant change.

A Soviet steel — Fe-13Cr-19Mn — approximates this limiting case

and therefore was studied. The actual chemical composition, by weight

percent, is: 13.40 Cr, 19.0 Mn, 0.82 Ni, 0.16 N, 0.055 C, 0.38 Si, 0.07

Cu, 0.020 P, 0.011 S, 0.10 V, 0.01 Mo, < 0.01 Ti , 0.0005 B, balance Fe.

Besides its scientific importance, this steel has technological applications. Intended for 1 iquefied-natural -gas tankage, it shows good low4

temperature fracture-mechanics parameters . The present communication

represents perhaps the first study of its physical properties.

Elastic constants were determined by measuring ultrasonic velocities

5 between 300 K and 76 K using equipment and procedures described elsewhere

82
The key experimental parameters include: orthogonal -prism specimen

1.3 cm by 1.3 cm by 1.6 cm; opposite faces flat and parallel within

3 ym; quartz transducers 1-cm diameter, x-cut and ac-cut, rough polished, gold electroplated two sides, 6 to 8 MHz fundamental frequencies;

bonding agent phenyl salicylate at room temperature, stopcock grease at

lower temperatures. Mass density determined pycnometrical ly was 7.81

3 g/cm . The raw material was- a 2.0-cm plate austenitized at 1 050°C for

1 h and air cooled.

Table I gives for selected temperatures the longitudinal and transverse sound velocities, v^ and v^, which Figs. 1 and 2 show as longitudinal and shear moduli

where p denotes mass density. Table I and Fig. 3 give four derived

[

1 ]

and

[ 2 ]

elastic constants^:

Shear modul us = G = C

t [3]

Bulk modulus = B

= [4]

Young's modulus = E = 9GB/(G + 3B) [5]

and

Poisson's ratio = v = (E/2G) - 1 [ 6 ]

83
Maximum experimental uncertainties are believed to be 1 percent for 6

and E and 1.5 percent for B and v.

All the elastic constants of this steel are anomalous near 274 K,

which corresponds presumably to its Neel temperature, considerably

higher than Tj^ = 215 K predicted by an empirical relationship developed

for Fe-Cr-Ni alloys^. Magnitudes of the elastic-constant changes agree

with previous results^. The overall results closely resemble those

reported for the highest-Mn alloy studied previously, Fe-1 8Cr-3Ni-l 3Mn

with a Ndel temperature of 191 K.

This steel fits the elastic-constant/temperature behavior pattern

3 established for Fe-Cr-Ni-Mn alloys . Reflecting its high Mn content,

its Neel temperature is the highest found thus far for these alloys.

The author acknowledges the technical assistance of M. Austin.

This study was supported in part by the National Science Foundation and

in part by the DoE Office of Fusion Energy.

84
H. M. Ledbetter: Mater. Sci. Eng. , 1977, vol . 29, p. 255-260.

E. W. Collings and H. M. Ledbetter: Phys. Lett., 1979, vol. 72A, p. 53-56

1 .

2 .

3. H. M. Ledbetter and E. W. Collings: in Metal Science of Stainless

Steels , TMS-AIME, New York, 1979.

4. R. L. Tobler, H. I. McHenry, and R. P. Reed: in Advances in Cryogenic

Engineering , Volume 24 , Plenum, New York, 1978, p. 560-572.

5. H. M. Ledbetter and D. T. Read: Metal 1 . Trans. A, 1977, vol. 8A, p. 1805-1808.

6 .

7.

L. D. Landau and E. M

1959, P. 13, P- 99.

L. A. A. Warnes and H

Lifshitz: Theory of Elasticity ,

W. King: Cryogenics , 1976, vol.

Pergamon, London,

16, p. 659-667.

85
Table I. Dynamic elastic properties of Fe-13Cr-19Mn steel at

selected temperatures, = longitudinal sound velocity, = transverse sound velocity, E = Young's modulus, G = shear

modulus, B = bulk modulus, and v = Poisson's ratio.

T

(K) O

>

''t cm/s)

E G

(10^^ N/m^)

B V

300 .5554 .3152 1 .959 .779 1.374 .263

290 .5554 .3155 1.961 .777 1 .373 .262

280 .5546 .3149 1 .954 .774 1.369 .262

270 .5495 .3098 1 .899 .749 1.359 .267

260 .5478 .3086 1 .885 .744 1 .352 .268

250 .5483 .3089 1 .889 .745 1 .354 .268

240 .5488 .3093 1.893 .747 1 .356 .267

230 .5495 .3099 1.900 .750 1 .358 .267

220 .5501 .3104 1.905 .752 1 .360 .267

210 .5508 .3108 1.911 .755 1 .363 .266

200 .5514 .3113 1 .916 .757 1 .365 .266

190 .5520 .3117 1 .921 .759 1 .368 .266

180 .5527 .3121 1.926 .761 1.371 .266

170 .5533 .3126 1.931 .763 1.373 .266

160 .5539 .3130 1.937 .765 1.375 .265

150 .5545 .3135 1.942 .767 1.378 .265

140 .5551 .3139 1.947 .770 1 .380 .265

130 .5556 .3144 1 .952 .772 1.381 .265

120 .5562 .3148 1.957 .774 1 .384 .264

no .5567 .3153 1.962 .776 1.385 .264

100 .5572 .3156 1.966 .778 1.387 .264

90 .5576 .3159 1 .969 .779 1 .389 .264

80 .5579 .3160 1.971 .780 1 .391 .264

86
1.01

LT5 CD

CNI

C_3

C_5

0.971 L

50 100 150 200 250 300

TEMPERATURE (K)

2 Fig. 1. Temperature dependence of the longitudinal modulus = pv^

for Fe-1 3Cr-l 9Mn . As described in the text, the anomalous behavior

is caused by a N^el (paramagnetic-anti ferromagnetic) transition.

87
C,(T)/Ct(295)

50 100 150 200 250 300

TEMPERATURE (K)

2 Fig. 2. Temperature dependence of the shear modulus G = pv^ for

Fe-13Cr-19Mn.

88
BULK

MODULUS

(10"N/m")

YOUNG'S

MODULUS

(10"N/m")

^ TEMPERATURE (K) Mg. J. Smoothed temperature dependence of the four elastic constants of

Fe-1 3Cr-l 9Mn : E - Young's modulus, G = shear modulus,

B - bulk modulus (reciprocal compressibility),

V = Poisson's ratio.

89

POISSON'S

RATIO

SHEAR

MODULUS

(lO^N/m^)

ROOM-TEriPERATURE ELASTIC CONSTANTS AND

LOW-TEMPERATURE SOUND VELOCITIES FOR SIX NITROGEN-STRENGTHENED

AUSTENITIC STAINLESS STEELS

National Bureau of Standards

91

ROOM-TEMPERATURE ELASTIC CONSTANTS AND LOW-TEMPERATURE SOUND VELOCITIES

FOR SIX NITROGEN-STRENGTHENED AUSTENITIC STAINLESS STEELS

H. M. Ledbetter

Fracture and Deformation Division

National Bureau of Standards

Boulder, Colorado 80303

Despite large chemical -composi tion differences, six ni trogen-al 1 oyed

stainless steels show only slight differences in their room-temperature

elastic constants, determined by ultrasonic-velocity measurements. At

low temperatures , magnetic transitions cause anomalous elastic-constant

behavior and significant differences in elastic constants. Nitrogen

content seems not to affect the elastic constants.

Key Words: Elastic constants; iron alloys; low temperatures ; magnetic

transitions; sound velocities; stainless steels.

93
Metallurgists began studying ni trogen-al 1 oyed stainless steels in

I 925 J Renewed interest in these alloys results from two causes:

concern over diminishing supplies of alloying elements such as nickel

and the unfulfilled dream of a high-strength face-centered-cubic steel.

Research on these alloys has been largely empirical and without

theoretical guidance. These studies, which focused mainly on mechanical

1 -3

strength properties have been reviewed by several authors.

One small but important step toward a basic understanding of a

material is knowledge of its elastic constants, which are among the most

fundamental mechanical properties. Effects of nitrogen on the elastic

constants of these alloys remain unstudied. Nor have such effects been

investigated for the relevant binary systems. Furthermore, the effects

of Cr, Ni, and Mn contents on the elastic constants of these stainless

steels remains unresolved. Another uncertainty is the occurrence of

low-temperature magnetic transitions, or possibly even structural

transitions, that manifest themselves as elastic-constant anomalies.

The presently reported study does not resolve these questions, but

contributes significantly toward their solution by providing accurate

measurements of basic properties of several alloys. The complete set of

elastic constants -- Young's modulus, shear modulus, bulk modulus, and

Poisson's ratio — were determined at ambient temperature for six

nitrogen-alloyed stainless steels representing a wide range of chemical

compositions. One of the elastic constants, the longitudinal modulus,

was determined during cooling to either liquid-nitrogen or liquid-helium

temperature.

4 Ledbetter and Col lings recently reported elastic constants of

three nitrogen-al loyed steels -- Fe-18Cr-3Ni-l 2Mn, Fe-21 Cr-6Ni-9Mn, and

94
Fe-22Cr-l 3Ni -5Mn -- between room temperature and liquid-nitrogen temperature. All these alloys undergo Neel (paramagnetic-antiferromagnetic)

transitions at a bel ow-ambient temperature, Tj^, that depends mainly on

the Mn content. All the elastic constants -- Young's modulus, shear

modulus, bulk modulus, Poisson's ratio, etc. -- behave anomalously at

temperatures near Tj,^. Higher Mn content causes a higher and a larger

elastic-constant change at Tj,^.

This communication reports results for three additional nitrogenalloyed steels -- Fe-18Cr-2Ni-13Mn, Fe-17Cr-9Ni-8Mn, and Fe-18Cr-8Ni

(AISI 304N), whose chemical compositions are given in Table I and

metallurgical characterizations in Table II. Note that all six steels

were in a sol ution-treated-and-quenched state; thus, most or all nitrogen was presumably in interstitial solid solution and not in precipitates. References 4 and references therein describe the experimental equipment and methods.

Table III gives the room- temperature sound-velocity and elasticconstant results. Excluding the Fe-1 7Cr-9Ni-8Mn alloy. Young's modulus

varies least (± 1.3 percent) and the bulk modulus varies most (± 2.4

percent). Considering the chemical-composition differences among the

alloys, these are surprisingly small variations. The slower shear-wave

velocity for Fe-1 7Cr-9Ni -8Mn , which results in lower Young's and shear

moduli and in higher bulk modulus and Poisson's ratio, is not understood. Strong texture resulting in higher elastic anisotropy may explain this result. But elastic anisotropy measured in Fe-18Cr-3Ni1 2Mn amounted to only 0.16 and 0.32 percent for the longitudinal and

5 shear velocities measured in three orthogonal directions. This alloy's

95
high Si content may cause its relatively low shear modulus, but apparently such factors remain unevaluated.

Longitudinal-modulus-versus-temperature curves shown in Fig. 1

reveal qualitatively similar behavior for all six alloys: a normal

linear increase upon cooling below room temperature, an anomalous

decrease through a minimum, and an increase to the zero-temperature

value. The longitudinal modulus reflects changes that occur either in

the bulk modulus or in the shear modulus. Neel (paramagneticantiferromagnetic) transition temperatures Measured previously^ for

g three alloys and computed empirically for one alloy are also indicated

as dashed vertical lines in Fig. 1. Clearly, the elastic-constant

anomalies relate to the magnetic transitions. Based on the elastic

data, Neel temperatures are approximately 171 K for Fe-18Cr-2Ni-l 3Mn and

43 K for Fe-17Cr-9Ni-8Mn.

In summary, six nitrogen-alloyed austenitic stainless steels, with

one exception, show similar room- temperature elastic constants. During

cooling all six steels undergo Neel (paramagnetic-antiferromagnetic)

transitions accompanied by anomalous elastic-constant/temperature

behavior. The transition temperature relates strongly to the Mn content. While further studies, with systematic nitrogen-content variations, are required, solid-solution nitrogen content seems not to

influence the elastic constants of these alloys. The Fe-18Cr-3Ni-12Mn

alloy was studied also in an annealed-and-furnace-cooled condition. The

1

2 elastic constants were: Young's modulus = 1.965’10 N/m , shear

modulus = 0.768-10^^ N/m^, bulk modulus = 1.489-10^^ N/m^, and Poisson's

ratio = 0.282. Thus, precipitation of nitrogen causes a one to three

percent lowering of both the dilatation-type and shear-type elastic

constants

96
The author acknowledges the contribution of alloy samples by Armco

Steel Corporation and experimental assistance by M. Austin. These

results emerged from a study sponsored by DoE Office of Fusion Energy.

97
REFERENCES

1. B. Gunia and G. R. Woodrow: J. Mater , 1970, Vol . 5, p. 413-430.

2. K. J. Irvine, T. Gladman, and F. B. Pickering: J. Iron Steel

Inst. , 1969, Vol. 207, p. 1017-1028.

3. A. Randak, W. Wessling, H. E. Beck, H. Steinmauer, and L. Faust:

Stahl und Eisen, 1971, Vol. 91, p. 1255-1270.

4. H. M. Ledbetter and E. W. Collings: in Metal Science of Stainless

Steel

, TMS-AIME, New York, 1979, p. 22-40.

5. H. M. Ledbetter and E. W. Collings: to be published.

6. L. A. A. Warnes and H. W. King: Cryogenics , 1976, Vol. 16, p. 659-667.

98
Table

I.

Chemical

compositions

(supplier's

analysis)

of

studied

alloys,

weight

percentage.

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oo • • 1 • • • • 0 0 0 0 0

00 LD LD LD 0 0 1 0 r— 0 OJ

oo 0 0 1 0 0 0 0 • • 1 • • • • 0 0 0 0 0 0

ro LD 1^ CTi

C\J OJ 1 OJ OJ OJ t—

Q_ 0 0 1 0 0 0 0 • • 1 • • • • 0 0 0 0 0 0 0 00 OJ 0 00

00 LO

O “ 0 0 0 0 0 0 0 0 0 0 0 0 0 • 00

•r“CO

ir> 0 >0 • OJ OJ OJ r— > CO CO OJ CO r— 1— fO CO • c 0 0 0 0 0 0 0 0 fO • 0 i-cu -0

00 OJ LO OJ LD CO 00 +-> • c

00 OJ 00 (Ti CTi 00 c >• fO c • O) LO

s: f— CO CO 00 00 1

C_) 1

r— r— r— • 0 4-> 0 s: c OJ • OJ •u to

t— OJ cn E c • LT) OJ OJ 00 CO LD Q. (T3 0 r“ • 0

CO CO LD OJ CO cri

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1- s: C_) fOcu 1

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S- s.. S- S- s_ s- s_

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r— r~ 1— OJ OJ f— r__ CO CO CO 1 1 1 1 1 1 1 O) r— f—

<u O) O) O) OJ cu O) 3 Ct c u. u_ u_ u_ U- Ll_ Lj_ fO JO u

99
Table

II.

Metallurgical

characterization

of

studied

alloys.

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100
Roomtemperature

sound

velocities

and

elastic

constants

of

six

nitrogen-strengthened

stainless

steels.

101
FIGURE CAPTION

2 Fig. 1. Temperature dependence of the longitudinal modulus pv^ for

six nitrogen-alloyed stainless steels. As described in

the text, the anomalies are due to Neel (paramagneticanti ferromagnetic ) transitions.

102
(

295 )

TEMPERATURE (K)

2 Fig. 1. Temperature dependence of the longitudinal modulus pv„ for

six nitrogen-alloyed stainless steels. As described Tn the text, the anomalies are due to Pleel (paramagneticanti fei^romagnetic) transitions

103

SOUND VELOCITIES AND ELASTIC CONSTANTS

OF AUSTENITIC STAINLESS STEELS 304, 310, AND 316

National Bureau of Standards

St*" ..K

/

Sound velocities and elastic constants of

austenitic stainless steels 304, 310, and 316

H. M. Ledbetter

Fracture and Deformation Division

National Bureau of Standards

Boulder, Colorado 80303

Room- temperature ultrasonic velocities, longitudinal and transverse,

were measured in three AISI 300-series austenitic stainless steels.

Elastic constants computed from these velocities show small but signi

ficant variations with composition.

Key Words

Bulk modulus; elastic constants; Poisson's ratio; shear

modulus; sound velocity; stainless steel; Young's modulus

107
Recently the author^ collaborated in studying the elastic-constant

variability in stainless steel 304. Twenty samples acquired randomly

were measured ul trasonical ly within an experimental uncertainty of 0.2

percent in the sound velocities. Results showed surprisingly small

elastic-constant variations: about 1 percent in the Young's and shear

moduli and about 1.5 percent in the bulk modulus and Poisson's ratio.

This included all possible variation sources: chemical composition,

thermomechanical processing, sample anisotropy, and experimental error.

This report describes a less-ambitious related study on two similar

materials: stainless steels 310 and 316 (three samples of each).

Experimentally, this study proceeded exactly as the previously reported

study^ on stainless-steel 304. On the basis of that study, several

possible sources of variation could be neglected: anisotropy; sampleto-sample, within lot; repeated measurements; measurement frequency.

Table 1 contains characterizations of the studied alloys.

Stainless steel 310 differs from 304 principally in containing considerably more nickel (by a factor of 2.5) and chromium (by a factor of

1.3) to stabilize the alloy against low- temperature strain-induced phase

transformati ons

Stainless steel 316 differs from 304 principally in containing an

additional two percent molybdenum to improve both corrosion resistance

and high-temperature mechanical properties. Also, it contains more

nickel by a factor of about 1.5, to compensate for the molybdenum, which

is a ferrite stabilizer.

Tables 2 and 3 show the results for stainless steels 310 and 316.

For comparison. Table 4 summarizes results ^ of the stainl ess-steel

108
304 study. In these tables, and denote the measured longitudinal

and transverse ultrasonic velocities, and p denotes the mass density.

Elastic constants relate to sound velocities through well-known relationshi ps

longitudinal modulus = = pv^

2 shear modul us = G = pv^

bulk modulus = B = p(v^ - v^)

Young's modulus = E = 3pv^ (v^ - ~ v^)

Poisson's ratio = v

= -^(v^ - 2v^)/(v^ - v^)

(

1 )

( 2 )

(3)

(4)

(5)

Since these three alloys consist principally of iron, chromium, and

nickel, some interesting observations arise concerning the relative

effects of chromium and nickel on the elastic constants. Contrary to

2 alloying effects in dilute alloys where increasing solute usually

decreases the elastic stiffnesses, in these alloys, chromium increases

all the elastic stiffnesses: B, E, and G. Nickel decreases both E

and G, as expected, but leaves B unaffected or slightly increased.

Chromium has little effect on Poisson's ratio, but nickel increases it

because

^ 4 , V ~ 9 i B

" G I

^

and the usual cancellation of the two terms does not occur.

The author is unaware of any systematic elastic-constant/composition studies of the basic Fe-18Cr-8Ni composition. Several binary

body-centered-cubiiron alloys, studied by Speich, Schwoeble, and

3 Leslie showed that chromium increased both B and G, while nickel

decreased both. Following previous studies, these authors factored the

109
observed alloying effects into electronic and volume contributions; they

concluded that Fe-Cr-Ni alloys show negligible electronic-term changes.

Further systematic studies of the elastic constants of alloys would

be useful in several ways: (1) effects of various alloying elements--

Cr, Ni, Mo, Mn-could be determined; (2) relationships of elastic

constants to theories of solid-solution strengthening and plastic deformation could be assessed; (3) fundamental properties would be provided

for testing various electron-theory models of transition metals; (4)

the important, unsolved problem of face-centered-cubic iron's elastic

constants might be elucidated.

ACKNOWLEDGMENT

This study was supported by the DoE Office of Fusion Energy and the

NBS Office of Standard Reference Data.

110
REFERENCES

1. H. M. Ledbetter, N. V. Frederick, and M. W. Austin: Elasticconstant variability in stainless steel 304, J . Appl . Phys

, 1980,

January issue.

2. C. Zener: Relation between residual strain energy and elastic

moduli, Acta Crystal logr. , 1949, 2, 163-166.

3. G. R. Speich, A. J. Schwoeble, and W. C. Leslie: Elastic constants

of binary iron-base alloys. Metal 1 . Trans

, 1 973 , Vol . 3, pp. 2031 -2037.

Ill
Table

1.

Chemical

compositions

(weight

percent)

and

hardness

(Rockwell

B).

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C\J

ZZZ • o o C\J CNJ C\J tn o o o • • • o o o

CD o 1

C\l CO

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CO CO

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CO CO • 1

c^ LO

C/1 • * « o o o

CO c 00

T~ • • * r—

CD CO

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CNJ •— LD VO

CO 00

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1

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s-o

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4-

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03

!Z03

03U

Q.

>3

CUo

cuCL

LT) Jm

112
Table 2. Sound velocities and elastic constants of stainless steel 310

at T = 295 K.

1 Sampl

2 3

Mean

Std.

Dev

Coeff.

of Var.

V^(cm/ys) 0.5765 0.5775 0.5749 0.5763 0.0013 0.2

V^(cm/ys) 0.3035 0.3049 0.3068 0.3051 0.0017 0.5

p(g/cm^) 7.850

Cj(GPa) 260.9 261 .8 259.5 260.7 1 .2 0.4

G(GPa) 72.3 73.0 73.9 73.1 0.8 1.1

B(GPa) 164.5 164.5 160.9 163.3 2.1 1 .3

E(GPa) 189.2 190.7 192.2 190.8 1.5 0.8

V 0.308 0.307 0.301 0.305 0.004 1 .2

113
Table 3. Sound velocities and elastic constants of stainless steel 316

at T = 295 K.

1 Sample

2 3

Mean

Std.

Dev

Coeff

of Var.

V^(cm/ys) 0.5712 0.5714 0.5648 0.5691 0.0038 0.7

V^(cm/ys) 0.3085 0.3100 0.3038 0.3074 0.0032 1 .0

p(g/cm^) 7.965 7.949 7.961 7.958 0.008 0.1

Cj(GPa) 259.9 259.5 254.0 257.7 3.3 1 .3

G(GPa) 75.8 76.4 73.5 75.2 1.5 2.1

B(GPa) 158.8 157.7 156.0 157.5 1 .4 0.9

E(GPa) 196.2 197.3 190.5 194.6 3.7 1 .9

V 0.294 0.291 0.296 0.294 0.003 0.9

114
Table 4, Sound velocities and elastic constants of stainless steel 304

at T = 295 K.

Mean Std. Dev. Coeff. of Var.

V^(cm/ys) 0.5759 0.0015 0.3

V^(cm/ys) 0.3134 0.0018 0.6

p(g/cm^) 7.880 0.045 0.6

Cn(GPa) 261 .3 2.0 0.8

G(GPa) 77.4 0.009 1.1

B(GPa) 158.2 0.022 1 .4

E(GPa) 99.6 0.017 0.9

V 0.290 0.004 1 .4

115

ELASTIC CONSTANTS OF STAINLESS STEELS

304, 310, AND 316 BETWEEN 4 AND 295 K

National Bureau of Standards

117

ELASTIC CONSTANTS OF STAINLESS STEELS 304, 310, AND 316 BETWEEN 4 AND 295 K*

H. M. Ledbetter

Fracture and Deformation Division

National Bureau of Standards

Boulder, Colorado 80303

For stainless steels 304, 310, and 316, longitudinal and transverse

ultrasonic velocities were measured by a pulse-echo method between 295

and 4 K. From these velocities were computed five elastic constants:

longitudinal modulus, shear modulus. Young's modulus, bulk modulus, and

Poisson's ratio. All three steels show low-temperature elastic-constant

anomal ies

Key Words: Bulk modulus; elastic constants; low-temperature; physical

properties; Poisson's ratio; shear modulus; sound velocity;

stainless steel; Young's modulus.

119
Several years ago^ the author collaborated in studying the lowtemperature elastic constants of four austenitic stainless steels: 304,

310, 316, and A286. That study indicated that all four steels show lowtemperature anomalies in dilatation-type elastic constants. But the sheartype elastic constants showed no measurable anomalies.

2 The study reported here, based on an improved measurement system,

reveals that the shear-type elastic constants are also anomalous at low

temperatures in two steels, 304 and 316. Steel 310 shows no anomaly in

its shear modulus. Besides better temperature sensitivity, the measurements

reported here have lower absolute uncertainty and represent averages of

several sources of samples, which were previously unavailable. The improved elastic-constant/temperature resolution elucidates the nature of

the relationship between elastic constants and 1 ow-temperature magnetic

phase transitions. Thus, the present study replaces that reported previously .

^

For brevity, experimental details are omitted here. Equipment and

2 techniques are described elsewhere. And the characterization of the

3 alloys are given in a companion publication.

Figures 1-5 show the elastic-constant/temperature results. And

elastic-constant values at selected temperatures are given in Table 1.

At temperatures exceeding about 100 K, all elastic constants of all

three alloys behave regularly: with increasing temperature, increasing

negative slope approaching linearity at higher temperatures.

With the exception of the shear modulus in stainless 310, all elastic

constants in all three alloys are anomalous at low temperatures . However,

in stainless 310, the anomalies in Young's modulus and Poisson's ratio are

rel atively smal 1

120
The shear-modulus anomaly occurs at the highest temperature in

stainless 304. The smaller, bulk-modulus anomaly occurs at the highest

temperature in stainless 310, where it occurs less abruptly over a wider

temperature region.

The "explanation" of these results lies in the magnetic transitions

4 5 in these alloys. As shown by Collings and Hart, ’ steels 304 and 316

undergo magnetic transitions near 50 K and 30 K, respectively ; but steel

310 shows no such transition, at least down to 5 K. Relationships between

magnetic transitions and elastic constants were reviewed by Ledbetter and

Col 1 i ngs

.

It should not escape attention that the elastic-constant changes and

magnetic transitions relate also to the crystal -structural stability of

the face-centered-cubic phase. Because of higher nickel content, steel 316

is more stable austenitical ly than steel 304. Steel 310 contains still

higher nickel, which is a strong austenite stabilizer, so that it resists

transition to the body-centered-cubic phase, caused either by cooling or

by mechanical deformation.

Since steel 310 shows no cooling-induced magnetic transition, the

shear-modulus and bulk-modulus anomalies must be related to separate

internal changes within the alloy. The paramagnetic-antiferromagnetic

transition, manifested as a peak in the magnetic-suscepti bi 1 i ty/temperature

curve, ^ is accompanied by a shear-modulus anomaly. The bulk-modulus

anomaly relates to another internal change at a higher temperature. This

different behavior of the shear and bulk elastic constants occurs in

6 7 several stainless-steel alloys and in invar alloys.

6-9

Comparing the present results with recent specific-heat studies

leads to a surprising conclusion. No specific-heat anomalies occur at
temperatures where the elastic constants change abruptly. Such anomalies

would be expected because the elastic constants determine the Debye temperature, which determines the lattice specific heat. This phenomenon was

discussed recently by Col lings and Ledbetter^*^ for an Fe-21 Cr-6Ni -9f1n

alloy, which exhibits a paramagnetic/spin-glass transition.

The 1 ow-temperature Young's moduli of two steels, 18Cr-8Ni and 18Cr-25Ni,

were reported recently by Verkin, Ilichev, and Klimenko. Their results

seem consistent with those of this study, although they invoke formation

of epsilon phase to partly explain the Young ' s-modul us/temperature behavior

of the less stable alloys. The reversi bi 1 i ty of our results on cooling

and heating suggests that epsilon phase is not a factor. Also, contrary

to the conclusions of Verkin et al .

,

we find no 1 ow-temperature anomalous

decrease in the Young's modulus of stable alloys.

This study was supported by the NBS Office of Standard Reference

Data and by the DoE Office of Fusion Energy.

122
1. H. M. Ledbetter, W. F. Weston, and E. R. Naimon, Low- temperature

elastic properties of four austenitic stainless steels, J. Appl

Phys. 3855-3860 (1975).

2. H. M, Ledbetter, N. V. Frederick, and M. W. Austin, J. Appl, Phys.,

(January 1980), forthcoming.

3. H. M. Ledbetter, Room-temperature sound velocities and elastic constants of austenitic stainless steels 304, 310, and 316, to be pub1 ished.

4. E. W. Collings and S. C. Hart, Low-temperature magnetic-susceptibility

and magnetization studies of some commercial austenitic stainless

steels. Cryogenics 1_9, 521 -530 (1979).

5. E. W. Collings and S. C. Hart, The magnetic character of austenitic

stainless steels, in Advances in Cryogenic Engineering , Volume 26

(Plenum, New York, 1980), forthcoming.

6. H. M. Ledbetter and E. W, Collings, Low-temperature magnetically

induced elastic-constant anomalies in three manganese stainless

steels, in Metal Science of Stainless Steels (Met. Soc. AIME,

New York, 1979).

7. G. Hausch, E, Torok, T. Mohri , and Y.“ Nakamura, Thermal expansion,

elastic properties, and the Mdssbauer effect of stainless invar,

J. Mag. Magnetic Mater. 1_0, 1 57-162 (1979).

8. J. M. Corsan and N, I. Mitchem, The specific heat of stainless

steels between 4 K and 300 K, in Proceedings of the Sixth Inter -

national Cryogenic Conference (IPC Science and Technology Press,

Surrey, 1976), pp. 342-344.

9. E. W. Collings, F. J. Jelenik, J. C. Ho, and M. P. Mathui', Magnetic

and thermal properties of stainless steel and inconel at cryogenic

temperatures , in Advances in Cryogenic Engineering , Volume 22 ,

(Plenum, New York, 1977), pp. 159-173.

123
10. J. C. Ho, G. B. King, and F. R. Fickett, Low temperature specific

heat of two stainless steels. Cryogenics 1_8, 296-298 (1978).

11. J. M. Corsan and N. I. Mitchem, The specific heat of fifteen stainless steels in the temperature range 4 K-30 K, Cryogenics 1^,

1 1 -1 6 (1 979).

12. E. W. Collings and H. M. Ledbetter, Sound-velocity anomalies near

the spin-glass transition in an austenitic stainless-steel alloy,

Phys. Lett. 7^_, 53-56 (1979).

13. B. I. Verkin, V. Ya. Ilichev, and I. M. Klimenko, The 1 ow- temperature

change of magnetic structure and plastic properties of Fe-Cr-Ni

alloys, in Advances in Cryogenic Engineering , Volume 26 (Plenum,

New York, 1980), forthcoming.

124
List of Figures

Fig. 1.

Fig. 2.

Fig. 3.

Fig. 4.

Fig. 5.

Longitudinal modulus, C^, of three stainless steels. =

2 pv^ , where p denotes mass density and denotes longitudinal

wave velocity.

2 Shear modulus, G, of three stainless steels. G = pv^ , where

p denotes mass density and v^ denotes transverse, or shear,

wave velocity.

4

Bulk modulus, B, of three stainless steels. B

=

~ 3 ^’

where denotes longitudinal modulus and G denotes shear

modul us

Young's modulus, E, of three stainless steels. E = 9GB/(3B+G),

where B and G denote bulk and shear moduli.

Poisson's ratio, v, of three stainless steels, v = (E/2G)-1,

where E and G denote Young's and shear moduli.

125
Fig. 1. Longitudinal modulus, C^, of three stainless steels.

pVp^, where p denotes mass density and denotes longitudinal

X/ wave velocity.

126
p denotes mass density and denotes transverse, or shear,

wave velocity.

127
170

165

160

155

150

304

100 200

TEMPERATURE (K)

300

k modulus, B, of three stainless steels. B = ^

G,

;re denotes longitudinal modulus and G denotes shear

lulus.

128
Fig. 4.

TEMPERATURE (K)

Young's modulus, E, of three stainless steels. E = 9GB/(3B+G),

where B and G denote bulk and shear moduli.

129
POISSON'S

RATIO

0.31

0.30 -

0.29

0.28

0.27

0 100 200

TEMPERATURE (K)

300

Fig. 5. Poisson's ratio, v, of three stainless steels, v = (E/2G)-1,

where E and G denote Young's and shear moduli.

130
TABLE 1. Elastic Constants of Three Stainless Steels at Selected Temperatures.

T V

(K) (GPa) (GPa) (GPa)

Stainless 304

295 199.6 77.4 158.2 0.290

280 200.9 77.9 158.3 0.289

260 202.5 78.7 158.8 0.288

240 204.1 79.3 159.4 0.287

220 205.6 80.0 159.7 0.285

200 207.1 80.6 160.0 0.284

180 208.6 81 .3 160.4 0.283

160 210.0 81 .9 160.6 0.282

140 211.3 82.5 161 .0 0.281

120 212.6 83.1 161 .1 0.280

100 213.7 83.5 161 .1 0.279

80 214.2 83.8 161 .2 0.279

60 214.3 83.9 160.5 0.278

40 211.0 82.5 159.1 0.279

20 210.1 82.1 158.4 0.279

5 210.1 82.1 158.3 0.279

Stainless 310

295 190.8 73.1

280 191 .9 73.6

260 193.4 74.2

240 194.9 74.8

220 196.2 75.4

200 197.6 76.0

180 199.0 76.6

160 200.4 77.1

140 201.6 77.7

120 202.8 78.2

100 203.9 78.8

80 204.8 79.1

60 205.7 79.4

40 206.2 79.7

20 206.4 79.8

5 206.4 79.9

163. 3 0, 305

163,,3 0,,304

164,.2 0,,304

164,.3 0,,302

164,.8 0,,302

165,.1 0,,301

165,,4 0,.300

166,,2 0,.299

166,.2 0,,298

166,.5 0,.297 166 .8 0 .296

166 .8 0,,295

166 .7 0,.294

166 .6 0 .294

166 .0 0 .293

165 .9 0 .293

Stainless 316

295 194.6 75.2

280 195.8 75.7

260 197.5 76.4

240 198.8 77.0

220 200.4 77.7

200 201.9 78.3

180 203.3 78.9

160 204.7 79.5

140 206.0 80.1

120 207.1 80.6

100 208.2 81.1

80 209.0 81 .4 60 209.4 81 .7 40 209.3 81 .7 20 207.8 81 .1 5 207.7 81 .0

157.5

158.0

158.3

158.8

159.1

159.5

159.7

160.0

160.3

160.5

160.6

160.5

160.2

159.6

159.0

158.9

0.294

0.293

0.292

0.291

0.290

0.289

0.288

0.287

0.286

0.285

0.284

0.283

0.282

0.281

0.282

0.282

131

YOUNG'S MODULUS OF A COPPER-STABILIZED

NIOBIUM-TITANIUM SUPERCONDUCTIVE HIRE

National Bureau of Standards

133

YOUNG'S MODULUS OF A COPPER-STABILIZED NIOBIUM-TITANIUM SUPERCONDUCTIVE

WIRE

H. M, Ledbetter

Fracture and_ Deformation Division

National Bureau of Standards

Boulder, Colorado 80303

J. C. Moulder

Thermo physical Properties Division

National Bureau of Standards

Boulder, Colorado 80303

M. W. Austin

Fracture and Deformation Division

National Bureau of Standards

Boulder, Colorado 80303

ABSTRACT

Young's modulus was determined for a 0.6-mm-diameter niobium-titanium

superconductive wire. Two methods were used; continuous-wave-resonance

and laser-pulse-excitation. Young's moduli were also determined for the

components -- copper and Nb-Ti — in both wire and bulk forms. Some

mechanical -deformation effects on Young's modulus were also measured.

From the components' elastic moduli, that of the composite was predicted

accurately by a simple rul e-of-mixtures relationship.

Key words: Copper; elastic constants; niobium-titanium; rule of

mixtures; sound velocity; superconductor; wire; Young's

modul us

135
I. INTRODUCTION

Most existing technological superconductors are wires. For example, in evaluating Nb^Sn's superconducting performance, Blum, Gregory,

and Coffey^ emphasized that its use depends on its "ability to be

handled and wound into superconducting devices... through proper design

2 and the use of rectangular wire..." Dew-Hughes reviewed practical

superconductive materials and concluded that "a superconductor is of no

use unless it can be manufactured as a continuous, flexible conductor...

this means the production of wire...." Several authors have discussed

3-5

the manufacture and nature of wire superconductors.

Elastic constants relate to several practical aspects of superconductive wire behavior in magnets and other devices -- load-deflection

behavior, thermoelastic stress, winding tension, prestress, magnetomechanical strain, residual stress -- as well as to other physical properties. Elastic constants relate in general to any stress-coupled force.

Depending on the symmetry of the stress and strain tensors, various

elastic constants -- for example. Young's modulus, shear modulus, bulk

modulus, Poisson's ratio, etc. -- become appropriate to a particular application.

Composite superconductive wires present particular difficulties to

the experimentalist measuring their elastic constants. In static methods,

stress must be low enough so that both matrix and fibers deform coherently and elastically. High-frequency ultrasonic methods may give an

apparent modulus too low because of dispersion or because of the diameterto-wavel ength ratio. Ultrasonic pulse methods often become difficult in

wires because of mode conversion caused by reflection of elastic waves

from the wire's surface. A composite-oscillator technique is impractical

136
because of the wire's small mass and small cross-sectional area. Thus,

as described below, unconventional experimental methods are required.

Especially for wires. Young's modulus, E, is the simplest, bestknown, and most useful elastic constant. Its utility arises from the

uniaxial stress, a^, that occurs along the wire axis, z. The axial

strain is e = a /E. For uniaxial stress the transverse strain is

A ^

e

= -va /E, where v denotes Poisson's ratio; when v is not isotropic,

A iC the second transverse strain, e^, differs from e^. In bending deformation, = a^/E = x/r, where x denotes a co-ordinate in the bending

direction perpendicular to z, and r denotes the radius of curvature of

the neutral surface near the origin. Velocities of extensional1 n compressional waves in rods and wires are given by v^ = (E/p) v.'here

p denotes mass density. As described below, this relationship provides

a simple, convenient, accurate method for measuring E.

The study reported here focused on Young's modulus of a single

superconductive composite wire: copper-stabilized Nb-Ti. The study had

three principal objectives: (1) determine the wire's Young's modulus,

which was uncertain; (2) determine the Young's moduli of the wire's components in both their annealed and deformed states, in both wire and

bulk forms; (3) evaluate critically the rul e-of-mixtures relationship

between component and composite Young's moduli.

137
II. EXPERIMENT

A. MATERIALS

Superconductor Wire

The composite superconductor wire (Fig. 1), obtained from a commercial

source, had 180 NbTi filaments in a high-conductivity, oxygen-free

copper matrix with final cross-sectional dimensions of 0.58 by 0.68 mm.

The filaments, 0.0322 mm diameter, were twisted 0.8 turns/cm. Their

chemical composition was 55 wt.pct. Nb and 45 wt.pct. Ti (38.7 at. pet. Nb,

61.3 at. pet. Ti). Based on a 1.8:1 CurNbTi volume ratio, computed mass

3 density was 7.921 g/cm based on 8.930 for copper and 6.104 for NbTi,

values measured hydrostatically on bulk samples. The wire was coated

with formvar insulation, a polyvinyl formal resin. Diamond-pyramid

hardness for a 2-g load was 98 ± 20; presumably this represents the

outer copper layer of the conductor. Residual resistance ratio was 70

and copper grain size was 0.02 mm. Being proprietary, the wire's thermomechanical treatment remains uncertain. But it should closely resemble

the treatment specified in the next paragraph for copper wire with the

addition of a final, large, cold reduction.

Copper Wire

High-conductivity, oxygen-free copper wire was obtained with the

following thermomechanical treatment: hot rolled at 1123 K to 0.6-cmdiameter rod, annealed at 1123 K for 0.5 h, water quenched, cold drawn

to 0.2-cm wire, annealed at 773 K for 0.5 h, water quenched, cold drawn

to 0.07-cm wire, annealed at 573 K for 1 h. Diamond-pyramid hardness

was 70 ± 10 for a 2-g load, residual -resistance ratio was 146, and

average grain diameter was 0.011 mm.

138
NbTi Bulk Alloy

Four-centimeter-diameter Nb-45Ti rod obtained commercially was

recrystal 1 ized at 1073 K and water quenched. Its oxygen content was

3

low, less than 1000 ppm. Mass density was 6.124 g/cm , room-temperature

electrical resistivity was 71 yfi*cm, diamond-pyramid hardness was

152± 6, and grain size was 0.04 mm and equiaxed.

The Nb-45Ti alloy was also given a thermomechanical treatment

resembling that of the superconductive wire; extrude from 4 cm to

1.55 cm diameter, anneal at 648 K for 48 h, draw to a final 1.0-cm

diameter. Hardness increased to 192 ± 5 in the deformed alloy, and

density decreased to 6.104. Most other properties remained unchanged,

although grain shape and grain size could not be determined unambiguously. Further properties of this alloy, particularly mechanical, were

c reported by Read, who detected a {110} extrusion texture typical of

body-centered-cubic alloys.

B. Methods

Two methods were used to determine Young's modulus of the wires:

continuous-wave-resonance and laser-pulse-excitation. While these

methods have not been applied extensively to° wires, the principles are

simple and the description here is brief. Young's modulus of bulk

materials was determined by either a pulse-echo-overlap method or by a

Marx-osci 1 1 ator method.

Continuous-wave-resonance

A length of wire between 5 and 10 cm was cemented between identical

PZT piezoelectric transducers 1.27-cm in diameter, whose fundamental

resonance frequencies were 200 kHz for longitudinal vibration.

139
Standing waves were established by sweeping the frequency of one transducer until a maximum voltage across the second transducer was detected

on an oscilloscope. Resonance corresponds to an integral number, n, of

half wavelengths in the wire. Thus, the velocity was obtained by:

Vg = fX = 2f£/n (1 )

where f denotes frequency, A wavelength, and I specimen length. For the

superconductor wire in resonance at 200 kHz, 10 half wavelengths equalled

a 9.55-cm specimen length. Thus, n and v^ were determined with little

difficulty by measuring, around the transducer's fundamental frequency,

several resonance frequencies of the wire, which were typically measurable

5 to 1 part in 10 or better. Further details of this method are described

7-8

by other authors.

Laser-pul se-excitat ion

Uide-spectrum acoustic pulses v/ere launched in wire specimens 3- to

5-cm long, by focusing the output of a Q-switched Nd-glass laser on one

end of the wire. The arrival of the acoustic pulse at the specimen's

opposite end was detected with an X-cut, 10 MHz-fundamental -frequency

quartz transducer, 0.32 cm in diameter, glued to the wire with phenyl

salicylate. The ends of the wires were ground flat by hand. The 14-ns

laser pulses were monitored with a portion of the beam reflected onto a

silicon photodiode; its output was summed el ectronical ly with the quartz

transducer output in the vertical amplifier of a storage oscilloscope,

triggered slightly before laser-firing by the Q-switch controller. This

produced an oscilloscope trace of two short pulses, one representing the

pulse launch time and one the pulse detection time. Transit times were

determined from photographs of the oscilloscope traces, which were

calibrated independently for each measurement with a time-mark generator.

Transit times determined using this technique were routinely reproducible

140
to within ± 1 percent; the overall accuracy in determining the extensional

velocity (calculated from the experimental velocity corrected for

dispersion) was determined by ijt inaccuracy, where £ denotes wire

length and t one-way transit time, is estimated to be ± 3 percent, owing

to the uncertainties associated with detecting the first motion of the

back surface and a smaller uncertainty in specimen length.

Pul se-echo-overl ap

A quartz piezoelectric crystal with a fundamental frequency near 10

MHz was glued with phenyl salicylate onto an approximately 1 cm-cube

specimen with flat and parallel surfaces. A 1-2 cycle acoustic pulse

was introduced into the specimen through the quartz transducer. The

pulse-echo train was detected by the same transducer and displayed on an

oscilloscope equipped with sweep delay and time-interval microprocessor

Sound velocities were determined by 2£/t, where £ denotes specimen

thickness and t round-trip transit time (time between echoes). Young's

modulus was computed from longitudinal and transverse sound velocities

combined with the mass density.

Marx oscillator

Cylindrical-rod specimens, 4.76-mm diameter, were cemented to a

matched pair of Y-cut, X-plated rectangul ar-rod quartz crystals with 50-

kHz fundamental frequencies. Using suitable electronics, this threecomponent (Marx) composite oscillator was swept in frequency until halfwave resonance occurred, which was detected by oscilloscope and measured

by frequency meter. The specimen's Young's modulus was determined from

the real tionshi p

E = 4pf^£^ (2)

where p = mass density, f = specimen frequency, and £ = specimen length.

141
III. RESULTS AND DISCUSSION

Table 1 gives results for sound velocities and Young's moduli in

the various materials. For wires, the extensional velocity, v^, is

obtained, which, when combined with the mass density, p, gives Young's

modul us

E = PVg^ (3)

For bulk specimens, longitudinal and transverse sound velocities, v^

and v^, were obtained, which give the extensional velocity by

2 . 2 , 2 4 2x,/ 2 2x

"e

' (''£ - 3^t

-

''t

>

A. NbTi bulk alloy

Young's modulus, 84.3 GPa for the annealed bulk NbTi alloy, agrees

approximately with previous studies. For example, from the resonance

9

study of Fedotov and Belousov one obtains an interpolated value of 79.3

GPa. Koch and Easton^ cite several other studies that give a similar

value for alloys with similar compositions. The compositional variation

of Young's modulus will be described extensively in a forthcoming publication by Ledbetter and Ceilings. As expected (since mechanical

deformation usually lowers elastic stiffness), the deformed alloy shov/ed

a lower Young's modulus, 79.2 GPa. The combined pulse-echo and Marxoscillator results for the deformed alloy show a significant texture,

6 amounting to stiffening along the rod axis, consistent with Read's

finding of a {110} rod texture.

B. Bulk copper

Copper's Young's modulus was determined recently in our laboratory

by both pulse and resonance techniques. These results, together with

several other results reported in the literature, v/ere described recently

by Ledbetter, who concluded that Young's modulus for annealed copper

is 128.7 ± 1.4 GPa.

142
C. NbTi Wire

Thus far, NbTi wire has not been measured in our laboratory. How1 2 ever, Schmidt reported a static stress-strain Young's modulus for Nb50 wt.pct.Ti wire, 0.26 mm in diameter, measured at liquid-helium temperature. Corrected for temperature and for composition, Schmidt's

result becomes 79.1 GPa. Based on the si ngl e-crystal studies of Reid,

1

Routbort, and Maynard for a Nb-43.6 wt.pct.Ti alloy, elastic anisotropy

is moderate for NbTi; the Zener anisotropy, A=2C^^/(C-| -j -C-J 2

) , is 1.76.

Compared with an isotropic Young's modulus of 87.4 GPa, this alloy has a

Young's modulus of 63.6 along [100], 91.8 along [110], 92.9 along

[111]. Thus, only a {100} wire texture would lower Young's modulus.

g Read's x-ray diffraction studies suggest a {110} texture, which would

slightly increase Young's modulus compared with the zero-texture case;

Schmidt's result supports Read's finding of a {110} rod texture. As for

bulk NbTi, mechanical deformation lowers the Young's modulus of the wire

by about six percent. This lowering could arise entirely from deformationinduced defects, and texture need not necessarily be a factor.

D. Copper wire

For annealed copper wire a Young's modulus of 126.2 GPa was ob14°

tained by cw-resonance ultrasonic-pulse methods at 200 kHz and a value

of 125.8 GPa by the laser-excitation method, giving a mean value of

126.0 GPa. This differs only one percent from the bul k-anneal ed-copper

value. As calculated from copper's singl e-crystal elastic constants.

Young's modulus would be 66.6 GPa for {100} texture, 76.8 GPa for {110},

191.4 GPa for {111}, and 128.7 GPa for the untextured, quasi-isotropic

case. Thus, the copper wire must be nearly, if not completely, texture

free.

143
E. NbTi :Cu wi re

For the composite superconductor wire, a Young's modulus of 107.8

GPa was determined by cw-resonance ultrasonic pulse^^ methods at 200 kHz

and 109.4 GPa by laser-pulse excitation; the mean value is 108.6 ± 1.1

GPa.

F. Rul e-of-mixture-predictions

If the strain along the fiber direction is the same in the fibers

and the matrix forming a composite, then it is simple to derive a ruleof-mixtures relationship for the effective Young's modulus in the fiber

direction

E = x^E^ + (l-Xf)E^ (5)

where f and m denote fiber and matrix and x denotes volume fraction.

This simple linear relationship is predicted by most elementary theories

that relate component and composite elastic properties. Actually, Eg.

(5) represents an upper bound; the lower bound being given by a similar

average for the elastic compliance, 1/E, rather than the elastic stiffness, E. This lower bound represents the Young's modulus perpendicular

to the fibers, where stress, not strain, is the same.

Using the bulk, annealed component values in Table I, Eg. (5)

predicts for the composite wire a Young's modulus of 109.3 GPa, within

one percent of the observed value. Using the component wire values

changes this prediction only slightly. Thus, in this case, a theory

more sophisticated than the rule-of-mixtures is not reguired to relate

the longitudinal Young's modulus to component Young's moduli.

144
IV. CONCLUDING REMARKS

From this study the follovying principal results and conclusions

emerge:

1. Copper wire processed thermomechanical ly to resemble the

superconductor matrix has the same Young's modulus as bulk,

annealed copper.

2. Deformed niobium-titanium wire has the same Young's modulus

as bulk, deformed niobium-titanium.

3. Young's modulus of an NbTiiCu wire superconductor can be

measured by various methods: ultrasonic-pulse, continuouswave-resonance, or laser-pulse-excitation. For this particular wire. Young's modulus was found to be 108.6 GPa.

4. Young's modulus for the composite superconductor wire can

be predicted reliably at room temperature using for the

Young's moduli of the components either bulk annealed values

or deformed-wire values. The predicted Young's modulus

was 109.3 GPa.

V. FUTURE STUDIES

Further studies should focus on the following problems:

1. Wires with other NbTi :Cu volume ratios and the applicability

of the rul e-of-mixtures to those wires.

2. Young's modulus perpendicul ar to the wire axis, either experimentally or theoretically.

3. Changes of Young's modulus of both components and superconductor with decreasing temperature. Applicability of ruleof-mixtures, especially at liquid-helium temperature.

4. Improved experimental methods for measuring a wire's elastic

properties.

145
5. Frequency spectrum of the detected acoustic signal for laserpulse excitation. The effect of diameter-to-wavelength ratio

for composite wires.

6. Component-composite elastic-property relationship for other

elastic constants such as the shear modulus and Poisson's

ratio.

7. Measurements of torsional modulus to provide a complete set of

elastic constants and to provide information on texture.

ACKNOWLEDGMENTS

This study was encouraged by Dr. J. W. Ekin of NBS, who supplied

the wires. Financial support was provided by the U.S. Naval Ship

Research and Development Center and by the DoE Office of Fusion Energy.

Dr. D. T. Read of NBS permitted use of his preliminary unpublished

studies of ultrasonic velocities in copper and superconducting wires,

Gary Maerz made several of the experimental measurements.

REFERENCES

1. P. Blum, E. Gregory, and D. L. Coffey, in Advances in Cryogenic Engineering, Vol . 22 (Plenum, New York, 1977), pp. 362-

369.

2. D. Dew-Hughes, in Superconducting Machines and Devices (Plenum,

New York, 1974), pp. 87-130.

3. D. Dew-Hughes, Mater. Sci . Eng. 1_, 2-29 (1966).

4. D. Dew-Hughes, Rep. Prog. Phys. 34, 821-73 (1971).

5. J. D. Livingston and H. W. Schadler, in Progress in Materials

Science, Vol. 12 , (Pergamon, Oxford, 1964), pp. 183.

146
6. D. T. Read, Cryogenics 1_8, 579-84 (1978).

7. D. I. Bolef and M. Menes, J. Appl . Phys. 1010-17, (1960).

8. D. I. Bolef and J. deKlerk, IEEE Trans. Ultrason. Eng. 1_0,

19-26 (1963).

9. S. G. Fedotov and P. K. Belousov, Fiz. Metal Metalloved 17 732

(1964).

10. C. C. Koch and D. S. Easton, Cryogenics ]

7_, 391 -413 (1977).

11. H. M. Ledbetter, J. Phys. D, forthcoming.

12. C. Schmidt, Appl. Phys. Lett. 463-5 (1976).

13. C. N. Reid, J. L. Routbort, and R. A. Maynard, J. Appl. Phys.

M, 1398-9 (1973).

14. D. T. Read, Fracture and Deformation Division, NBS, Boulder,

Colorado, personal communication.

147
Table I. Sound velocities and Young's moduli for a superconducting

wire and its components

Material Density

(g/cm^)

Extensional

Sound Velocity

(cm/ps)

Young '

Modul us (GPa)

Cu, bulk, annealed 8.93 0.380 128.7

Cu, wire, annealed^ 8.93 0.376 126.0

NbTiiCu, wire, deformed^ 7.92 0.370 108.6

NbTi:Cu, wire, rul e-of-mixtures 7.92 0.371 109.3

NbTi, wire, deformed^ 6.10 0.360 79.1

NbTi, bulk, annealed 6.12 0.371 84.3

NbTi, bulk, deformed 6.10 0.360 79.2

^Includes unpublished pulse-echo resultsJ^

^Extrapolated from results in Ref. 12.

148
FIGURE CAPTION

Fig. 1. Photomicrograph of transverse section of superconducting v/ire.

NbTi filaments, 180 in number, constitute 5/14 of the volume.

The wire's dimensions are 0.58 by 0.68 mm.

149
Fig. 1. Photomicrograph of transverse section of Nb-Ti superconductive

wire. MbTi filaments, 180 in number, constitute 5/14 of the

volume. The wire's dimensions are 0.58 by 0.68 mm.

150
Welding

151

WELDMENTS FOR LIQUID HELIUM SERVICE

National Bareau of Standards

m

153

WELDMENTS FOR LIQUID HELIUM SERVICE

H. I. McHenry and T. A. Whipple

National Bureau of Standards

Boulder, Colorado 80303

The program to evaluate weldments for structural applications in

superconducting magnet systems [1, 2] was continued in FY79. The

objectives for the program are;

1. To investigate the metallurgical factors which affect

the toughness of stainless steel weld metals at cryogenic

temperatures.

2. To contribute to the development of improved filler

metals for welding stainless steels for liquid helium

servi ce

,

3. To evaluate the strength and toughness of stainless steel

and aluminum weldments at 4 K.

During the three years of the program, the first objective has remained

unchanged, and objectives 2 and 3 have been generalized to cover the

expanding scope of the program. Research started this year on the

evaluation of promising filler metals for improved-toughness stainless

steel weldments (but not improved strength), on the evaluation of

aluminum alloy weldments, and in the development of improved procedures

for measuring fracture toughness at 4 K. Work continued on metallurgical

investigations and toughness tests on stainless steel weldments supplied

by various contributors to the DoE/OFE superconducting magnet programs.

This report summarizes the results obtained in FY79 and briefly outlines

the plans for FY80.

Stainless Steel Weldments

Previous work [1, 2] has identified ferrite content, nitrogen

content, and sensitization as causes for toughness deterioration in

stainless steel weldments at cryogenic temperatures. Most of the

investigations were done on AWS 31 6L welds deposited by the shielded

155
metal-arc (SMA) process in 25.4-mm thick plates. This year, numerous

toughness tests were conducted on 31 6L SMA welds which were ferritefree, and contained low levels of N and C. The welds were supplied by

participants in the Large Coil Project and the Mirror Fusion Test Facility

Project. The results, presented later in this report by Whipple,

McHenry, and Read, indicate that wide variations in fracture toughness

occur at 4 K even when ferrite is absent, nitrogen is constant, and

sensitization is nil. The toughness at 4 K correlates well with grain

size, indicating that process variables and welding procedures can have

a marked influence on toughness.

A series of 308L weldments were supplied for liquid helium testing

in support of the LCP magnet design. Welds were deposited by the SMA,

gas metal arc (6MA), and flux core metal-arc (FCMA) welding processes.

The welds contained between 3 and 8% ferrite. The strength and fracture

toughness values measured at 4 K are summarized in Tables 1 and 2.

Brittle behavior was observed in the FCMA welds where the fracture

toughness was measured to be50 MPav^ using valid plane strain procedures

per ASTM E 399. The low toughness was attributed to high nitrogen in

the welds; the FCMA welding process is self-shielded (i.e., no inert gas

coverage is used) and tends to pick up nitrogen from the atmosphere.

A welding process study for 316L stainless steel weld metal was

conducted by Kotecki . The results of this investigation are summarized

later in this report. The purpose of this investigation was to assess

the effect of welding process and, to a lesser extent, welding parameters on the cryogenic toughness of AWS 316L weld metal at two nominal

ferrite levels, 0 and 5%. Welds were deposited in 25.4-mm thick plate

156
by the gas tungsten-arc (GTA), GMA, and submerged arc (SA) processes.

Fissure bend tests, room temperature tensile tests, and Charpy tests at

77 to 173 K were conducted, and test plates for evaluation at 4 K were

provided to NBS. The results indicate that GTA and GMA welds have

acceptable toughness at 77 K by ASME standards, but the SA welds have

unacceptable toughness. Further tests using the J-integral method will

be conducted by NBS in FY80.

Two experimental stainless steel weld metals were evaluated at 4 K.

The properties are summarized in Tables 3 and 4. A 1 6Cr-l 6Ni -2Mo alloy

was deposited by the FCMA process. The weld metal had relatively low

strength, 606 MPa, and excellent fracture toughness, 239 MPav^m. The flux

core wire used to deposit the 1 6Cr-l 6Ni-2Mo weld contained a special

flux formulated to result in a low nitrogen deposit under open-arc

conditions (i.e., without inert gas coverage). The outstanding toughness of this alloy combined with the improved deposition rates possible

with the FCMA process makes this alloy an excellent candidate for

further developmental work. The second experimental alloy was the 18Cr16Ni-9Mn alloy, which was previously evaluated at 77 K [2]. The 18 Cr16Ni-9Mn weld deposited by the GMA process had relatively high tensile

strength, 1076 MPa, but marginal fracture toughness, 120 MPa/m. At the

present time, no further work is planned on this alloy because of the

low fracture toughness and the excessive mi crofi ssuring , which lead to

an almost totally intergranular fracture.

Two SA weldments have been evaluated at 4 K. The results are given

in Tables 5 and 6. The SA weldments were supplied by the MHD project

being conducted by MIT. The two weldments contained 4.7 and 8.0%

ferrite and had fracture toughness values of 136 and 99 MPa/ii, respectively.

157
Aluminum Alloy Weldments

A program to evaluate the tensile and notched tensile properties of

several aluminum alloy 5083 welds at 4 K was conducted by Kelsey and

Mueller [3]. The idea was to have several fabricators and suppliers

prepare test welds in 5-cm thick 5083-0 plate using 5183 filler wire.

The equipment, procedures, and filler wire source were selected by each

of the six participating companies. All welds were deposited by the GMA

process. The welded plates were tested to determine the variability of

weld properties at 4 K. The results of this program are presented later

in this report by Kelsey. Portions of each test weldment will be sent

to NBS for J-integral fracture toughness testing at 4 K. The results

will be correlated with the notched tensile test results to determine if

the empirical correlations used for aluminum alloys at room temperature

and 77 K are val id at 4 K.

Test-Method Development

Fracture toughness testing is the principal method for evaluating

weldments at 4 K. The J-integral test procedure used is expensive, time

consuming, and difficult. A simplified procedure is needed for research

purposes and also for weld qualification testing by fabricators and

independent laboratories. Two test-method-development activities are

discussed later in this report. A computer-interactive J-integral test

procedure was developed to improve the efficiency of J-integral testing

at NBS-Boulder. This method is described by Read, Tobler, and Reed [4].

A simple test method for estimating Jj^ was developed at NBS by Begley

and evaluated for liquid helium service by Whipple and McHenry. The

method appears to be suitable to base metals, but further testing,

scheduled for FY80, is needed to determine its applicability to weldments.

158
Future Work

The main emphasis in the FY80 program will be to start work on the

development of materials and procedures to improve the toughness of

stainless steel welds deposited by high deposition rate processes.

Initial work will be directed toward defining the metallurgical characteristics necessary for good toughness in 1 arge-grai n-si ze welds.

Fracture toughness testing of stainless steel weldments will continue in

support of the LCP program, the weld process selection study, the 16Cr16Ni-2Mo filler metal development, and a cooperative program with an

Austrian filler metal supplier on ferrite free SMA welds. In addition,

work will continue on the evaluation of the 5083 aluminum alloy weldments and on development of the estimation procedure.

REFERENCES

1. H. I. McHenry, "Evaluation of Stainless Steel Weld Metals at Cryogenic

Temperatures," in Materials Studies for Magnetic Fusion Energy

Applications at Low Temperatures

, NBSIR 78-884 (1978), pp. 159-

166.

2. H. I. McHenry, D. T. Read, and P. A. Steinmeyer, "Evaluation of

Stainless Steel Weld Metals at Cryogenic Temperatures," in Material

Studies for Magnetic Fusion Energy Applications at Low Temperatures-II , NBSIR 79-1609 (1979), pp. 299-308.

3. R. A. Kelsey and L. N. Mueller, "Tensile and Notch-Tensile Properties

at Room Temperature and 4 K of Welds in Aluminum Alloy 5083-0

Plate," in Material Studies for Magnetic Fusion Energy Appi i cations

at Low Temperatures-II , NBSIR 80-1627 (1 980), pp.

4. D. T. Read, R. L. Tobler, and R. P. Reed, "J-Integral Methods and

Measurements for AISI 304 Stainless Steels at Low Temperatures"

submitted for publication (1 980).

159
Tensile

Properties

of

308L

Weldments

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K

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160

Magne-Gage

measurements
Fracture

Toughness

and

Fatigue

Properties

of

308L

Weldments

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161

3)

da/dN

in

mm/cycle;

aK

in

MPa»4ii’
Tensile

Properties

of

Experimental

Filler

Metals

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162
Fracture

Toughness

and

Fatigue

Properties

of

Experimental

Filler

Metals

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163
Tensile

Properties

of

316L

SA

Weldments

at4

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164
TABLE 6

Fracture Toughness and Fatigue Properties of 316L SA Weldments at 4 K

Percent

Ferrite 2

'^ic

kj/m (in-lb/in )

Kic (0)

MPa/m (ksi/fn)

da/dN = Cak"^

^ • -2 metric units

^ (English units )

4.7 84

7178 (443)

138

126

132 (121)

--

' 8.0 47

43

40

43 (247)

103

98

96

99 (90)

1 .22 X 10‘]p 3

(7.13 X 10^ ^)

1) Magne-Gage measurements

2) da/dN in mm/cycle; aK in MPa/m

3) da/dn in in/cycle; aK in ksi/i*n

I

I

165

FRACTURE BEHAVIOR OF FERRITE-FREE

STAINLESS STEEL WELDS IN LIQUID HELIUM*

T. A. Whipple, H. I. McHenry, and D. T. Read

National Bureau of Standards

Boulder, CO 80303

Work supported by the U.S. Department of Energy, Office of Fusion Energy.

167

ABSTRACT

Most research to date concerning the cryogenic toughness of stainless steel weldments has concentrated on the effects of delta-ferrite

content and nitrogen concentration. It has been shown that an increase

in either of these leads to reduced cryogenic toughness. This paper

reports that a wide variation in fracture toughness occurs in 31 6L

weldments that contain no delta ferrite and only small differences in

nitrogen concentration. Therefore, welding parameters and minor compositional variability must also have significant influences. Optical

microscopy revealed a strong correlation between cryogenic fracture

toughness and fusion-zone grain width, which is controlled by welding

parameters and weld-metal composition. An increase in grain size reduces

the toughness. Microfissuring, which is a problem in low-ferrite welds,

did not have an effect on the measured fracture toughness, but in one

case it severely reduced the tensile ductility. Fatigue crack growth

rate tests showed little variation in the fatigue properties of 31 6L

weldments, regardless of the variations in other properties.

169
INTRODUCTION

The fabrication of the load-bearing members of large superconducting

magnet support structures requires the welding of thick sections. These

weldments must be capable of withstanding high stresses at liquid

helium temperatures. Current welding practice yields stainless steel

weldments with sufficient strength; however, the high stresses, thick

sections, and low temperatures increase the possibility of brittle

fracture [1]. Therefore, the fracture toughness and fatigue crack

growth rate of the weldments are important design parameters. Previous

work on the cryogenic toughness of austenitic stainless steel weldments

has generally centered on the effects of two important variables: delta

ferrite content and nitrogen concentration [2,3,4].

Read, McHenry, Steinmeyer, and Thomas [2] showed that large decreases

in both fracture toughness and Charpy impact energy were related to

increasing delta-ferrite content at 76 K. At 4 K, a strong decrease in

fracture toughness with increasing delta-ferrite content was also observed.

As the del ta-ferri te content was increased from 0.13% to 10.1%, the

Charpy impact energy decreased from 35 to 12 J, and the decreased

2 from 135 to 41 kJ/m at 4 K. These results demonstrate convincingly the

strong detrimental effect of delta ferrite on low temperature fracture

toughness.

Szumachowski and Reid have also investigated the role of ferrite in

reducing the low temperature toughness of stainless steel weldments [3].

Using 77 K Charpy impact tests, they found that the toughness not only

increased with a decrease in the actual ferrite content, but continued

to increase as the calculated ferrite number became negative. Also, lowcarbon content and the use of lime-coated electrodes improved the low

170
temperature toughness of stainless steel weldments. In a continuation

of this study, Szumachowski and Reid found that an increase in the

nitrogen content of the weldment led to a decrease in the 77 K Charpy

impact energy [4]. This is consistent with the results of Tobler and

Reed [5], who reported that nitrogen was an extremely effective strengthening agent for 304 stainless steel base metals at cryogenic temperatureSj but the increase in strength with increasing nitrogen led to a

decrease in the toughness.

In this paper, recent data on the tensile and fracture properties

of stainless steel weldments at 4 K are presented, and factors other

than ferrite and nitrogen that appear to affect these properties are

di scussed.

MATERIALS

The data reported here have, for the most part, been obtained from

weldments that were produced as part of a test program or are qualification weldments associated with large superconducting magnet programs

sponsored by the Department of Energy. The magnet programs include the

Large Coil Project (LCP) at Oak Ridge National Laboratory and the

Mirror Fusion Test Facility (MFTF) at Lawrence Livermore Laboratory.

The weldments discussed here are shiel ded-metal -arc (SMA) weldments made

with special low-ferrite AWS E316L-15 electrodes. These weldments

contain 0% delta ferrite (by Magne-Gage measurements) and were, with the

exception of one vertical weld, produced in the flat position.

PROCEDURES

The weldments were evaluated using tensile, fatigue crack growth

rate and fracture toughness tests. All tests were conducted in liquid

hel ium at 4 K.

171
Tensile Tests

Tensile tests were conducted at 4 K using 2.5-cm (1-in) gage length

and 0.64-cm (0.25-in) diameter specimens. The specimens were oriented

along the length of the weldment such that the entire sample was welded

metal. The test procedures and the cryostat for the low temperature

tests have been described by Reed [6].

Fracture Toughness Tests

Fracture toughness tests were conducted at 4 K using the ASTM E399

compact specimen with a 5.1-cm (2-in) width and 2.5-cm (1-in) thickness.

The specimens were generally oriented such that loading was perpendicular to the weldment, and crack extension was along the weldment

centerline. The single-specimen J-integral procedure [7] was used to

determine the value of J at the onset of crack extension. The plane

strain fracture toughness, Kj^(J), was determined from using the

relationship proposed by Begley and Landes [8];

where E is Young's modulus and v is Poisson's ratio. Tests were conducted in an electrohydraul ic fatigue machine equipped with the cryostat

described by Fowlkes and Tobler [9].

In the single-specimen J-integral technique, the specimen is precracked in tension-tension fatigue to a crack length of 3 cm (1.2 in).

Subsequently, the specimen is loaded monotonical ly at a load-linedisplacement rate of 50 um/s or less under displacement control to a

point within the elastic region. The curve of load vs. load-line

displacement is recorded on a second x-y plotter with appropriate

amplification and signal conditioning.

172
A signal proportional to the change in load is electronically substracted

from the signal; the proportionality factor is adjusted so that the unloading compliance before crack extension is displayed as a vertical

line on the graph of load change against load-line displacement change.

The specimen is incrementally extended, with several stops for recording

the unloading compliance.

Typically, ten to twenty unloading compliances are recorded. The

magnitude of the decrease in load during the unloadings is restricted to

10% of the total load. J-integral values are calculated from the

load vs. load-line displacement recorded at each unloading point using

the Merkl e-Corten relation [10]. Crack extension values Aa, are calculated from all unloading compliance curves except the first one, which

is used to evaluate the compliance at Aa = o. The (J vs. Aa) data are

then plotted, and is evaluated graphically as the intersection of a

curve through the data, with the theoretical blunting line given by J

=

2oAa, where a is the arithmetic mean of the yield strength and the

ultimate tensile strength.

This method evolved into the following procedure, which was also

used in this study: Load and displacement signals are converted to

digital form, and a laboratory minicomputer processes the data and

plots each (Aa, J) datum point immediately after the associated unloading

is completed [11]. The Aa values obtained from the compliance data

generally agreed within 10% of those measured after the test. Scatter

in the J-resistance curves led to uncertainties of about ± 12% in the

J. values.

Ic
Fatigue Crack Growth Rate Tests

Fatigue crack growth rate tests were conducted at 4 K on compact

specimens identical to those used for fracture toughness testing. The

tests were conducted in the same testing machine and cryostat described

above. The specimens were sinusoidally loaded in tension with an

amplitude of 90% of the peak load and at a cyclic rate of 20 Hz. At

regular intervals, the compliance of the specimen was measured on an x-y

plotter. The intervals between compliance measurements were shortened

as the crack growth rate increased so that roughly the same amount of

crack extention occurred between each compliance measurements were shortened

growth per cycle, da/dN, was calculated from the compliance measurements,

and the stress intensity change, AK, was calcualted from the crack

length and loads. These data are plotted according to the relationship:

da/dN = CAK*^

where C and n are constants.

RESULTS AND DISCUSSION

Tables 1 and 2 contain a summary of the data obtained during this

investigation on 316L SMA weldments, which contained no delta ferrite.

Fracture Toughness Properties

The results of Read et al . [2] showed that, in a controlled series

of 316L SMA weldments, an increase in the delta-ferrite content resulted

in a decrease in the fracture toughness. In that series of weldments,

the chromium and nickel contents were varied while the compositions of

minor alloying elements and the welding parameters were held constant.

This approach resulted in a very good description of the effect of

174
del ta-ferrite content. In the present study, a number of different 316L

SMA weldments with 0% ferrite and similar nitrogen contents were tested.

Although all these weldments had a similar composition, they were

produced with different welding parameters and in different plate

thicknesses. Figure 1 shows the (J) data of Read et al . along with

the fracture toughness data for 316L SMA weldments obtained during this

study. There is a larger variation in fracture toughness for weldments

that contain no ferrite than that obtained for the controlled series of

weldments containing 0 to 10% ferrite. Therefore, the fracture toughness

is not only a function of the ferrite and nitrogen contents, but is also

strongly dependent on other factors, such as welding parameters and

other compositional variations.

Attempts to correlate the variation in fracture toughness with

other mechanical properties, such as tensile properties or Charpy VNotch (CVN) toughness, have been unsuccessful. The range of yield

strengths for the 0% ferrite 316L weldments is only 772 to 882 MPa

(112 to 128 ksi); this small variation in yield strengths can not

account for the large difference in fracture toughness values of 95 to

200 MPa»/fn~ (86 to 182 ksi/Tn). Also, 77 K CVN toughness does not vary

greatly within this group of weldments and does not correlate with the

fracture toughness. Figure 2 shows the strong correlation found between

the fracture toughness, Kj^(J), and the grain size of the fusion zone.

The fusion zone of stainless steel weldments consists of large columnar

grains that display epitaxial growth, that is, the grains are continuous

between passes (see Figure 3, a three-dimensional micrograph of the

weldment structure). When a compact tension specimen is loaded, the

plane that experiences the highest strain is the plane parallel to the

175
plate surface. The structure of this plane is essentially equiaxed;

therefore, the important grain size for deformation is the columnar

grain width, this is the grain size plotted in Figure 2.

This dependence of the toughness on the grain size leads to some

interesting possibilities concerning the effects of specimen orientation

on the mechanical properties of the weldment. For example, for cracks

propagating in the through-thickness direction, the columnar grains are

oriented with the length direction in the direction of crack propagation. This has the effect of increasing the average grain size in the

plane of highest strain. One weldment, No. 7, was tested in both

orientations, the crack propagating down the centerline and in the

through-thickness direction. The results are shown in Table 1. Although the load-deflection curve of the through-thickness sample was

still character! Stic of a tough material, the value was only

107MPai/m (97 ksi/Tn) compared with 200 MPavirT (183 ksi/Tn) for the

standard orientation. This is only the result of one test and is not

conclusive. However, it does point out that a design assumption of

isotropic weldment properties may not be valid, and with regard to

normal stress states, a crack propagating in the through-thickness

direction is usually very important. This aspect of weldment properties

deserves further investigation.

It has also been observed that there is a directional dependence to

the plastic deformation of stainless steel weldments. An example of the

directional dependence is seen in the necked down section of round,

longitudinal tensile specimens, in which the fracture surface is always

elliptical in shape. The long axes of the grains are aligned with the

short axis of the ellipse, showing that plastic deformation occurs more

176
easily in that direction. An example of the deformation following the

columnar grains in a compact tension specimen is shown in Figure 4.

This specimen is oriented such that the crack is propagated in the

through-thickness direction. Notice that the deformed region is not

symmetric around the crack tip as it would be in an equiaxed structure,

but follows the length of the columnar grains.

Effect of Microfissures

Given consistent weldin-g conditions, a decrease in the deltaferrite content of the weldment results in an increased cryogenic

fracture toughness [2] However, the production of very low ferrite or

fully austenitic weldments can be difficult because these weldments are

apt to form microfissures. The subject of microfissuring has been

extensively studied [12-16], particularly with regard to the effects of

del ta-ferrite content and chemical composition. It is generally agreed

that microfissuring is the result of a liquation process, where a lowmelting temperature grain-boundary constituent is remelted by the

subsequent welding pass. The thermal stresses cause the fissure to

open.

The exact cause of microfi ssuring is not yet well established;

nevertheless, it is well known that microfissuring can be eliminated by

adjusting the composition to obtain 5 to 10% delta ferrite. The amount

of delta-ferrite necessary depends on the chemical composition and the

welding process and parameters used. To date, the technology to produce

consistently microfissure-free , fully austenitic weldments is not

available. Therefore, the effect of these defects on mechanical properties must be considered.

177
From their studies of the effects of microf issures on mechanical

properties, Honeycombe and Gooch [13] concluded that microfissures are

unlikely to have a significant effect on the tensile or Charpy impact

properties of the weldment. It was observed that microfissures may

extend by ductile tearing, but have no propensity for unstable failure.

They did find that microfi ssures can have an effect on the fatigue life,

if the microfissures are the most significant defects present (i.e.,

if the weldment reinforcement is removed and there are no other defects).

In this study, no apparent effect of microfissures on the fracture

toughness was found. For weldment No. 8 which contained significant

microfissuring , the fracture toughness value was consistent with other

similar weldments that did not contain microfissures. This observation

is not surprising, considering the orientation of the microf issures with

respect to the crack plane. The microfissures, being perpendicular to

the crack, act as obstacles to crack propagation and cause the crack to

branch or change plane.

In weldment No. 8, the microfissures were generally oriented

transverse to the welding direction, that is, they were perpendicular to

the tensile axis of all weld-metal tensile samples. In this orientation,

microfissures reduce the load-bearing area of tensile samples. This did

not appear to affect the yield strength, but it did severely reduce the

ultimate strength and ductility measurements. The total elongation

values for the three tensile samples tested were 2.2, 8.6, and 13.8%.

This variation in total elongation appears to relate directly to the

size of the microfissures on the fracture surface -- the largest microfissure was observed on the sample with the lowest elongation. The

magnitude of this effect is related to the area of the microfissure

178
relative to the entire cross-sectional area, and since the size of the

microfissures is generally small, a larger sample size should reduce the

effect. Therefore, if the entire cross section of the weld is considered,

the effect of microfissures on tensile ductility will probably be small.

They, will only be significant if the loading is in the longitudinal

direction of the weld.

Figures 5 and 6 are scanning electron microscope (SEM) fractographs

of the fracture surface of weldment No. 8. Figure 5 shows a grain

boundary that was completely open prior to testing; note that the

liquated surface shows no evidence of deformation. Although the microfissure appears to have opened up during testing, it does not appear to

have propagated. Figure 6 is an example of a microfissure in which the

grain boundary did not separate completely during welding. It is evident

that there were small ligaments of solid material left in the grain

boundary; these ligaments failed in a ductile manner during the test.

Fatigue crack growth rate tests were also conducted on the weldment

with microfissures. No effect of microfissures was observed; this is

probably also due to the orientation of the microfissures with respect

to the crack plane, as discussed above.

Fatigue Crack Growth

The results of the fatigue crack growth rate tests are given in

Table 2 and shown graphically in Figure 7 for the low-ferrite 316L SMA

weldments. Figure 8 combines the data for all the weldments tested and

shows only minor variations in the fatigue crack growth rate, in spite

of the fact that these weldments represent fracture toughness variations

from 132 to 200 MPa/iT” (120 to 182 ksi/Tn). In all cases, the fatigue

crack growth rate properties of the weld metal are better than those for

179
304 N stainless steel base metal [17], which is representative of the

base metals used for magnet casings.

SUMMARY

The tensile fracture toughness and fatigue crack growth properties

of eight 316L stainless steel SMA weldments were evaluated in liquid

helium at 4 K. The tensile strength and fatigue crack growth rates were

similar for all eight weldments. However, the fracture toughness

varied from 95 to 200 MPa /m. The toughness increased with decreasing

fusion-zone grain width. Microfissuring was observed in several weldments but with one exception, did not appear to influence the properties

measured, in the excepted weldmdnt, the tensile strength and ductility

were reduced due to microfissures in the test section.

180
REFERENCES

1. H. I. McHenry and R. P. Reed, "Structural Alloys for Superconducting Magnets in Fusion Energy System," Nucl.Eng. Des

,

forthcoming

2. D. T. Read, H. I. McHenry, P. A. Steinmeyer, and R. D. Thomas, Jr.,

"Metallurgical Factors Affecting the Toughness of 316L SMA Weldments at Cryogenic Temperatures ," Material Studies for Magentic Fusion

Energy Applications at Low Temperatures-II , NBSIR 79-1 609, National

Bureau of Standards, Boulder, CO, pp. 313-352.

3. E. R. Szumachowski and H. F. Reid, "Cryogenic Toughness of SMA

Austenitic Stainless Steel Weld Metals: Part I - Role of Ferrite,"

Weld. J. , Weld. Res. Suppl . 57 , 325-s-333-s (1 978).

4. E. R. Szumachowski and H. F. Reid, "Cryogenic Toughenss of SMA

Austenitic Stainless Steel Weld Metals: Part II - Role of Nitrogen,"

Weld. J. , Weld. Res. Suppl

58, 34-s-44-s (1 979).

5. R. L. Tobler and R. P. Reed, " Intersti tual Carbon and Nitrogen Effects

on the Tensile Properties and Fracture Mechanics Parameters of an AISI 304

Type Austenitic Stainless Steel", Material Studies for Magnetic

Fusion Energy Applications at Low Temperatures-III . , NBSIR 80-1627,

National Bureau of Standards, Boulder, Co, pp. 15-48.

6. R. P. Reed, "A Cryostat for Tensile Tests in the Temperature Range

300 to 4 K," in Advances in Cryogenic Engineering, Vol . 7, Plenum

Press, New York (1962) pp. 448-454.

7. G. A. Clarke, W. R. Andrews, P. C. Paris, and D. W. Schmidt, in

Mechanics of Crack Growth , Proceedings of the 1974 National Symposium on Fracture Mechanics, ASTM STP 590, American Society for

Testing and Materials, Philadelphia, PA (1976), pp. 27-42.

131
8.

J. A. Begley and J. D. Landes, in Fracture Toughness , Proceedings

of the 1971 National Symposium on Fracture Mechanics, Part II, ASTM

STP, 514 American Society for Testing and Materials, Philadelphia,

PA (1972), pp. 1-20.

9. C. W. Fowlkes and R. L. Tobler, "Fracture Testing and Results for a

Ti 6A1-4V Alloy at Liquid Helium Temperatures , " Eng . Fract. Mech. ,

8, 487-500 (1976).

10. J. G. Merkle and H. T. Corten, "A J-Integral Analysis for the Compact Specimen, Considering Axial Force as well as Bending Effects,"

Journal of Pressure Vessel Technology, Trans. ASME, 96 , 1-7 (1974).

11. J. A. Joyce and J. P. Gudas, Computer Interactive J ^^ Testing of

Naval Alloys ASTM STP 668, American Society of Testing and Materials,

Philadelphia, PA (1979) pp. 451-468.

12. C. D. Lundin and D. F. Spend, "The Nature and Morphology of Fissures

in Austenitic Stainless Steel Weld Metals," Weld. J. Weld. Res. Suppl

356-S-367-S (1976).

13. J. Honeycombe and T. G. Gooch, "The Effect of Microcracks upon the

Mechanical Properties of Austenitic Stainless Steel Weld Metals,"

Report No. M/69/72, The Welding Institute, Cambridge, MA (1972).

14. S. Kanazawa, K. Yamato, T. Tekeda, and K. Hashimoto, "Study of

Reheat Cracking in Weldment (Report I)," Trans. Jpn. Weld. Soc. ,

7, No. 1, 49-57 (1976).

15. J. Honeycombe and T. G. Gooch, "The Effect of Compositional and

Process Variables on Microcracking in Fully Austenitic Stainless

Steel and Weld Metal," Report No. M/74/73, The Welding Institute,

Cambridge, MA (1973).

182
16. J. A. Brooks and F. J. Lambert Jr., "The Effects of Phosphorus,

Sulfur and Ferrite Content on Weld Cracking of Type 309 Stainless

Steel," Weld. J. Weld. Res. Suppl . 57 , 1 39-s-143-s (1 978).

17. D. T. Read and R. P. Reed, "Toughness, Fatigue Crack Growth, and

Tensile Properties of Three Nitrogen-Strengthened Stainless Steels

at Cryogenic Temperatures," Materials Studies for Magnetic Fusion

Energy Applications at Low Temperatures-I , NBSIR 78-884, National

Bureau of Standards, Bou-lder, CO, (1 978), pp. 91-154.

183
Tensile

Properties

of

316L

Stainless

Steel

SMA

Weldments

at4

K.

-MO

13 -a

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184

845

(123)

*Magne-Gage

measurements

indicate

0%

ferrite.
Toughness

and

Fatigue

Properties

of

316L

Stainless

Steel

SMA

Weldments

at4

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List of Figures

Fig. 2

Fig. 3

Fig. 4

Fig. 5

Fig. 6

Fig. 7

Fig. 8

Critical stress intensity factor, (J) as a function of

del ta-ferrite content at 4 K, showing the wide variation in

fracture toughness that has been found for 316L SMA weldments

with 07o ferrite.

The relationship between fusion-zone grain width and 4 K

fracture toughness of 316L SMA weldments with 0% ferrite.

Composite three-dimensional photomicrograph of the fusion zone

of a 0% ferrite 31 6L SMA weldment.

Macrograph of the deformed region around the crack tip of

a compact specimen oriented such that the crack is propagated

in the through-thickness direction, showing that plastic

strain is transmitted down the length of the columnar grains.

SEM fractograph of a microfissure on the fracture surface of

weldment No. 8, (a) 150 X (b) 1000 X.

SEM fractograph of a grain boundary, which was only partially

separated prior to testing, on the fracture surface of weldment

No. 8, (a) 50 X (b) 2000 X.

The 4 K fatigue crack growth rate data for 31 6L SMA weldments

(Nos . 4 to 8)

The 4 K fatigue crack growth rate data for all 316L SMA weldments

tested, compared with 304N base-metal data from Read and Reed [17].

186
CRITICAL

STRESS

INTENSITY

K.^iJ)

(MPa/m)

PERCENT FERRITE

Fig. 1. Critical stress intensity factor, K. (J), as a function of

del ta-ferrite content at 4 K, showing the wide variation in

fracture toughness that has been found for 31 6L SMA weldments

with 0% ferrite.

137
250

200

150

100

50

T T T T

4 K

o

O (MFTF)

I 1_

.2 0.4

_L_

0.6

AVERAGE GRAIN WIDTH (mm)

e relationship between fusion-zone grain width and 4 K acture toughness of 316L SMA weldments with 0 % ferrite.

laa
j

1

L - Longitudinal or welding direction

T - Long transverse or width direction

S - Short transverse or weld thickness direction

Fig. 3. Composite three-dimensional photomicrograph of the fusion zone

of a 0 % ferrite 316L SMA weldment.

iRq
Fig. 4. Macrograph of the deformed region around the crack tip of a compact specimen oriented, such that the crack is propagated

in thru thickness direction showing that plastic strain

is transmitted down the length of the columnar grains.

190
!

1

SEM Fractograph of a microfissure on the fracture

surface of weld No. 8, (a) 150 X (b) 1000 X. Fig. 5.
Fig. 6. SEM Fractograph of a grain boundary, which was only partially

separated prior to testing, on the fracture surface of weld

No. 8 , (a) 50 X (b) 2000 X.

192
(aiD/CD/ujui) 3iva HlMOaO XDVMD

193

The

4

K

fatigue

crack

growth

rate

data

for

31

6L

SMA

weldments
CRACK

GROWTH

RATE

da/dN

(mm/cycle)

STRESS INTENSITY CHANGE AK (MPa/m)

Fig. 8. The 4 K fatigue crack growth

tested, compared with 304N ba

rate data for all 31 6L SMA weldments

se-metal data from Read and Reed [17].

1 94
I WELD PROCESS STUDY FOR 31 6L STAINLESS STEEL

t'k,' WELD METAL AT CRYOGENIC TEMPERATURES

ft

‘ '

Tel edyne-Mc Kay

195

Weld Process Study for 31 6L Stainless Steel

Weld Metal at Cryogenic Temperatures

D . J . Ko tec k

Taledyne - HcKay

Scope

In response to an inquiry from the National Bureau of Standards, Thermophysical

Properties Division, Boulder, Colorado, Teledyne McKay proposed to produce a series of type 316L stainless steel weldments. Six of these weldments were to

be made in 1-inch thick type 304L base plate which would then be returned to NBS

for their tests. In addition, under each of the same six welding conditions,

al 1 -weld-metal specimens and fissure bend specimens were to be prepared and tested

at Teledyne McKay.

The purpose of this program was to assess the effect of welding process choice and

to a lesser extent welding parameters choice on cryogenic toughness of type 316L

weld metal at two nominal ferrite levels - approximately 5FN and approximately 0 FN.

This report describes the welding procedures used to produce the weldments for NBS

evaluation and describes the procedures and test results of the al 1 -wel d-metal

experiments at Teledyne McKay.

Filler Metals

Due to present steel mill melting practices, type 31 6L bare wire is fairly easy to

obtain with ferrite calculated by the Delong diagram as about 5FN. But unless high

nitrogen is acceptable, type 3161 bare wire calculated to have 0 FN is rare. From

work done previously at Teledyne McKay (1>2), it is apparent that high nitrogen has

a detrimental effect on cryogenic toughness. In order to begin with a low nitrogen

wire of 0 FN calculated, tubular wire of special composition was produced for this

work.

From the available inventory of type 31 6L wire heats, a heat produced by Car Tech

(ht. ^23847) was selected. The mill analysis of this heat was reported as follows:

Mn Si Cr Ni_ P_ L-

^

021 1.58 .50 18.61 12.79 2.17 .019 .002 .044 .08

Cu

.11

The calculated ferrite by the Delong Diagram is 5.2FN using the mill analysis

197
A tubular wire was produced aiming deliberately to exceed the 11.0-14.0 percent

nickel range for type 31 6L wire in AWS A5.9, so that 0 FN might be obtained. The

aim wire composition was as follows:

Mn Si lil Mo

.02 2.3 0.4 18.5 15.5 2.3 .02

The calculated ferrite by the Delong Diagram for this composition would be 0.2FN.

Approximately 220 lbs of this wire were produced. Samples were taken from four

locations within this lot (lot #9230-497) approximately 60 lbs apart and GTA melts

were prepared and checked for nickel, chromium and molybdenum homogeneity. Satisfactory homogeneity was indicated and the wire was accepted for use in the study.

Sample Preparation

Both the bare wire and the tubular wire were used at 3/32-inch diameter for all

test welds.

For the welds in the 1-inch 304L plates to be tested at NBS, the basic joint preparation was a 90° included angle double vee which was symmetrical about the plate

midthickness. No root opening was used and, in the case of GTA welding, no land was used either. For GMA and SA welding, a land of 1/8-inch + 1/16 was used to

prevent burn-through. The joints were fit up and restrained by welding each plate

to a stainless steel bar at each end to prevent the root from opening ahead of the

arc. After the first side of the joint was welded, the root was backgouged with

the air-carbon arc process to sound metal, then lightly ground to remove dross

from gouging before welding the second side. After completion of welding, the

surfaces were ground smoothe, then radiographed. If a defect was found, it was removed by air-carbon arc gouging and repaired before shipping the weldment to

NBS.

For the al 1 -wel d-metal studies of mechanical properties, all weld joints were in

3/4-inch thick plate with a 45° included angle single vee joint preparation, a 1/2-inch root opening and a 3/8-inch backup bar. The plate and backup bar used

were mild steel. The joint faces and the backup bar were covered by 2 layers of

stainless weld metal from the wire under test by the welding process under test

before assembling the joint preparation. (This method of producing an all -weldmetal joint is given in AWS A5. 22-74. AWS A5.9-77 which applies to the wires

under test has no joint preparation.. For each wire and welding procedure, an

al 1 -weld-metal plate was made under the same conditions (except joint preparation)

as for the 304L joints shipped to NBS. From each al 1 -wel d-metal joint, a longitudinal

tensile specimen and 12 full-size Charpy vee notch specimens were machined according

to the locations given in AWS A5.5-69. None of the specifications applying to stainless filler metals - A5.4, A5.9 or A5.22 - specifies Charpy vee notch specimens.

198
For each wire and each welding procedure, a chemistry/ferri te pad was produced

by building up 8 layers (1 pass per layer) of weld metal. The weld was water

quenched after each pass except the last two passes. The next-to-last pass was allowed to air cool for 1 minute before quenching. The last pass was air cooled

to room temperature. The top surface was then lightly draw-filed as recommended

in AWS A5.4-78. Six readings along the weld centerline were taken with a MagneGage calibrated according to AWS A4.2-74 and averaged to obtain the FN of the

sample. A light milling cut was then taken from the surface to obtain chips

for nitrogen and oxygen analysis. ,The surface remaining was then lightly ground

before running remaining chemical analysis with an emission spectrometer.

For each wire and each welding procedure, a fissure bend specimen was prepared.

This consisted of 2 layers of weld beads deposited as stringer beads lengthwise

on 3/8x3xl0-inch type 304 plate. The second layer consisted of 4 passes. After

each pass, the specimen was air cooled to black heat, then water quenched. The

specimen was restrained flat by welding it to a 1-inch thick plate before beginning

the stringer beads. After completion of welding, the top layer of stringer beads

was surface ground, removing just enough metal to provide a clean surface at each

of the three stringer bead interfaces where fissures are generally found. The

ground surface was then examined using fluorescent dye penetrant and any indications

in the center 4 inches of the length were counted. The specimen was then bent

around a 1-1/2-inch diameter to a 120° angle and again examined with fluorescent

dye penetrant. The indications were counted in the center 4 inches as before.

The fissure bend test used herein is basically as described by Lundin et al(3).

Welding Procedures

Three welding procedures were used with the 316L bare wire. The first was a manual GTA procedure using 1/8-inch tungsten (2% thoria) at 230-240 amps DCSP

with argon shielding. Cold wire addition was made by feeding the wire through a hand-held GMA gun which did not supply current. The on-off switch on the gun

provided a convenient means of adding filler at will without interrupting the

welding to change filler wires as would be necessary with straight cut lengths

of filler wire. For the 1-inch 304L plate provided to NBS, no filler was used

in the root pass only, and a total of 13 passes were used to complete the joint.

Each layer of weld was a single pass except the top layers which v;ere two passes

each. For the al 1 -weld-metal joint in 3/4-inch plate, filler was used for all

passes. Fourteen passes, two per layer, were used to complete the joint. These

weldments are labeled 9230-508-1.

The GMA procedure for the 31 6L bare wire consisted of semi-automatic welding at

400 amps, 28-30 volts DCRP in 99% Argon - 1% oxygen gas at 40 cfh with 3/4-inch

electrode extension. The 1-inch 304L plate required 4 single-pass layers (2 per

side) and 2 two-pass layers (1 per side) for a total of 8 passes. The all-weldmetal joint used a single-pass root, 4 two-pass layers and a single three pass

layer for a total of 12 passes. These weldments are labeled 9230-508-2.

199
The SA procedure for the 31 6L bare wire consisted of machine welding at 300

amperes, 29 volts DCRP, 16 inches per minute travel, 1 inch electrical extension

with Lincoln STIOO flux. For the 1-inch 304L plate, a single pass root and second

layer on each side was followed by two two-pass layers and one three-pass layer

for a total of 18 passes. For the 3/4-inch al 1 -wel d-metal plate a single pass

root was followed by five 2-pass layers and a single 5-pass cover layer to a. total of 16 passes. These weldments are labeled 9230-508-3.

Three welding procedures were used with the modified 316L tubular wire also.

The first of these consisted of semi-automatic GMA welding at 400 amps, 28-30 volts

DCRP in 99% Argon - 1% oxygen gas at 40 cfh with 3/4-inch electrode extension. The

1-inch 304L plate required a single-pass root on each side and two 2-pass layers

per side to complete for a total of 10 passes. The al 1 -wel d-metal plate required

a single pass root, three 2-pass layers and one 3-pass layer for a total of 10

passes. These weldments are labeled 9230-497-5.

The second procedure with the tubular wire was machine SA welding at 300 amps,

29 volts DCRP, 16 inches per minute travel, 1-inch electrical extension with

Lincoln STIOO flux. The 1-inch 304L plate used a single-pass root on each side, one 2-pass layer on one side and two 2-pass layers on the second side, and one 3-pass layer on each side for a total of 14 passes. These weldments are labeled

9230-497-6.

The third procedure with the tubular wire was machine SA welding at 400 amps,

32 volts DCRP, 10 inches per minute travel, 1-inch electrical extension with

Lincoln STIOO flux. The 1-inch 304L plate required a single pass root on each

side, one single pass layer on one side and two single-pass layers on the other,

and a two-pass layer at the top of each side for a total of nine passes. The

al 1 -weld-metal plate used a single pass root, three 2-pass layers and one 3-pass

layer for a total of 10 passes. These weldments are labeled 9230-497-7.

For each 1-inch 304L plate and for each al 1 -wel d-metal plate, a maximum interpass

temperature of 300F was employed.

Results

The radiographs of all of the 1-inch 304L plate weldments were free of defects

except the high heat input SA weld with the tubular wire. This showed slag entrapment. The defect was removed by air-carbon arc gouging followed by light grinding

to remove dross. A repair was then made by machine SA at the original welding

conditions. Radiography confirmed that a sound repair was made. All of the

1-inch 304L plate weldments have been shipped to NBS for evaluation.

Al 1 -weld-metal deposit ferrite and chemical composition were obtained and are given

in Table 1. In comparing the SA deposits to the GMA and GTA deposits, several

composition changes will be noted. These are to be expected with the Lincoln

STIOO flux. There is about a 0.5% increase in Mn and Si with this flux and

nearly a 1.0% increase in Cr. An increase in phosphorus, oxygen and carbon is

also apparent in the SA deposits.

200
In comparing the tubular wire deposits with the bare wire deposits, it will be

noted that the bare wire deposits are higher in several residual elements - nitrogen, cobalt, vanadium and copper. On the other hand, the tubular wire

deposits are higher in titanium. This is as expected. The melting of stainless

steel heats in the steel mill involves use of scrap which introduces residual

elements. On the other hand, the tubular wire contains a small deliberate

addition of titanium for improved GMA operation.

It is worthy of comment that both 'the bare wire and tubular wire deposits contain

appreciable variations in nickel from test to test - nearly a 1% difference

between high and low values in both cases. These differences are not due to

interactions with the SA flux. Their origin is unexplained.

The measured FN values for the bare wire deposits are higher than for the

corresponding tubular deposits (GMA and SA). The FN values calculated using

the DeLong diagram do not in all cases agree well with the measured FN values.

The gain in Si and Cr with SA welding is not offset by gains in C and Mn, so that FN's in SA welds tend to run higher than in GTA or GMA welds.

The fissure bend test results are given in Table 2. Despite the low FN values

measured on some of the welds, the fissure counts in all cases are quite low.

All of the fissures noted are along and perpendicular to the interface between

two weld passes in the weld metal which is reheated by a subsequent weld pass.

The al 1 -weld-metal tensile and impact values obtained are given in Table 3.

The room temperature tensile and yield strengths for the bare wire run appreciably

higher than for the tubular wire. This is presumably due to the higher ferrite

and nitrogen of the bare wire deposits.

The Charpy vee notch impact results show, for the wires tested, significantly

better toughness for the GMA deposits than for GTA or SA deposits. The better

results for GMA than for GTA seem to agree with our internal observation in

stainless covered electrodes that larger diameter rods (bigger beads) give

better toughness at cryogenic temperatures than smaller diameter rods (smaller

beads). Note also that high heat input SA does better than normal SA with the

tubular wire. It is not clear why SA should do noticeably poorer in cryogenic

toughness than the other processes - higher silicon and oxygen as shown in Table 1 may be part of the problem with SA. However, other tests we have run with fully

austenitic weld metal of 1 6Cr-l 6Ni-2Mo have shown equal toughness (30 mils lateral

expansion at -320F) for SA and GMA deposits. It should be noted that, while the

high heat input SA weld did fairly well in toughness, the welder who made these

weldments felt this was not a good way to weld - it is too easy to trap slag

because the bead is so large it tends to roll over on the edges.

Based on our tests with covered electrodes as reported in References 1 and 2,

better toughness would have been expected with the low FN, low nitrogen tubular

201
wire than with the bare wire for the same welding process. This is certainly

not true for the GMA and SA welds tested herein. Since the oxygen and silicon

are virtually identical for GMA and normal heat input SA welds with the two wires,

the only available explanation for this behavior would seem to involve residual

titanium in the tubular wire deposits or lack of other residual elements.

The final evaluation of the test weldments must await the fracture toughness 'data

to be developed at NBS on the 1-inch type 304L plates.

References

1. E. R. Szumachowski and H. F. Reid, "Cryogenic Toughness of SMA Austenitic

Stainless Steel Weld Metals; Part 1 - Role of Ferrite", Welding Journal , V57, Nil (November, 1978) , pp 325s-333s.

2. E. R. Szumachowski and H. F. Reid, "Cryogenic Toughness of SMA Austenitic

Stainless Steel Weld Metals: Part 2 - Role of Nitrogen, " Welding Journal , V58, N2 (February, 1979), pp. 34s-44s.

3. C. D. Lundin, W. T. DeLong and D. F. Spond, "Ferrite-Fissuring Relationship

in Austenitic Stainless Steel Weld Metals", Wei ding Journal , V54, N8 (August, 1975),

pp. 241s-246s.

202
Table 1 A11-We1d Metal Composition and Ferri te

31 6L Bare 316L Tubular

SA high input

GTA GMA SA GMA SA input

9230-508-1 9230-508-2 9230-508-3 9230-497-5 9230-497-6 9230-497-7

C .017 .017 .023 .014 .016 .017

Mn 1.39 1 .40 1.97 2.17 2.50 2.43

Si .45 .47 .84 .46 1 .03 .84

Cr 18.02 17.95 18.95 18.68 19.61 19.07

Ni 11.89 11.76 12.75 15.15 15.32 14.59

Mo 2.12 2.12 2.08 2.31 2.21 2.22

S .003 .003 .005 .013 .011 .011

P .020 .020 .031 .010 .021 .021

N .069 .069 .065 .030 .016 .018

0 .053 .066 .110 .072 .106 .087

Co .092 .090 .095 .012 .025 .020

V .063 .063 .06 .013 .017 .016

Cb .02 .02 .02 .02 .02 .02

Ti .021 .022 .024 .15 .092 .104

Cu .089 .086 .10 .037 .043 .042

FN measured 1.5 3.0 7.4 0.9 2.2 1 .8 FN Calculated

by Diagram 4.4 4.1 5.7 0.8 6.8 5.2

Table 2 Fissure Bend Test Results

316L Bare 316L Tubular

Fissure Size

GTA GMA

9230-508-1 9230-508-2

SA GMA

9230-508-3 9230-497-5

SA

9230-497-6

SA high heat

i nput

9230-497-7

Before Bending

poi nt

up to 1/16"

1/16-1/8" over 1/8"

0

0

0

0

0

0

0

0

7 0 0

0 0 0

0 0 0

0 0 0

0

0

0

0

After Bending

point 0

up to 1/16" 0

1/16-1/8" 0 over 1/8" 0

0

0

0

0

160

0

0 1 0

1 0 0

0 0 0

0 0 0

203
Table 3 Tensile and Impact Test Results

316L Bare 316L Tubular

SA high

GTA GMA SA GMA SA heat input

9230--508-1 9230- 508-2 9230- 508-3 9230-497 -5 9230- 497 -6 9230-497-7

Tensil

Strength, ksi 87.31 80.34 85.14 72.97 78.57 80.53

Yield

Strength, ksi 62.40 52.04 63.92 48.06 56.08 46.67

% Elongation 32 45 36 30 32 34

% Reduction in Area 54.6 54.9 53.6 33.6 41 .2 37.8

Charpy vee notch

ft-lbs (mils)

@-150 78.8 (38) 83.9 (43) 36.9 (15) 39.0 (4o;1 19.5 (10]1 41.2 (24)

73.8 (29) 77.4 (41) 36.2 (19) 49.2 (4i:1 18.1 (15]1 32.5 (24)

73.8 (43 91 .8 (54 34.7 (14 40.5 27)1 14.5 i 10]1 26.0 (15

avg. 75.5 (36) 84.4 (46) 35.9 (16) 42.9 (36]1 17.4 1(l2j1 33.2 (2l)

@-200 57.1 (39) 89.7 (57) 31.1 (14) 34.0 (28) 13.7 ( r,1 28.2 1(19)

73.8 (45) 70.1 (40) 32.5 (18) 32.5 (22) 13.7 ( 7]1 28.2 1(24)

65.1 34 81.0 ( 43 ;I 32.5 16) 37.6 32 14.5 ( 8]1 33.3 1 21 avg. 65.3 (39) 80.3 ( 47 ;

1 32.0 (16) 34.7 (27) 14.0 ( 7]1 29.2 1(21)

@-250F 54.2 (27) 70.9 (44]1 27.5 OD 23.9 (14) 13.0 ( 6)1 27.5 (:i6)

62.9 (27) 60.7 (30]1 25.3 (10) 33.3 (18) 13.7 ( 7)> 23.1 (:ii)

49.2 28 72.3 44 1 28.2 pi 29.6 20) 13.7 ( 7 23.1 (;i4

avg. 55.4 (27) 68.0 ( 39 )

1 27.0 (11) 28.9 (17) 13.5 ( 7 ) 24.6 (. 14 )

@-320F 41.2 (12) 52.1 (26)1 21.0 ( 2) 25.3 (18) 11.6 ( 1) 31.8 (17)

47.7 (24) 62.2 (31)i 21.0 ( 2) 20.2 (15) 10.1 ( 3) 31.1 (15)

39.8 (18) 61.5 (39) 22.4 ( 2) 23.9 (17) 13.7 ( 5) 23.1 (12)

avg. 42.9 (18) 58.6 (32) 21.5 ( 2) 23.1 (17) 11.8 ( 3) 28.7 (15)

204
THE COMPUTER-AIDED J-INTEGRAL TEST FACILITY* AT NBS

National Bureau of Standards

> (

•

r205 K,

?!.

THE COMPUTER-AIDED J- INTEGRAL TEST FACILITY AT NBS

D. T. Read

ABSTRACT

The computer-aided J-integral testing facility that has been implemented at NBS-Boulder over the past year is described briefly.

INTRODUCTION

Acceptance of the parameter as a measure of the fracture resistance of metals has become wide'spread in recent years. A large volume

of J-integral fracture toughness data on a variety of candidate materials

for low temperature applications has been accumulated. J-integral

testing has the advantage of being applicable to smaller specimens (up

to 100 times smaller) and tougher materials then the plane strain fracture

test (Kj^ test). A disadvantage of the J-integral technique has been

its experimental difficulty. Previous versions of this test have required

approximately five specimens and lengthy data reduction procedures [!]•

Recently developed single-specimen J-integral test techniques [2] have

eliminated the multiple specimen requirement, but the data reduction

procedure was still lengthy. In the computer-aided J-integral test,

most of the data-reduction procedure is automated; the apparatus produces

a plot of J-integral value against crack extension from which Jj^ can be

quickly and simply extracted.

THE J. MEASUREMENT PROBLEM

Ic The measurement problem in Jj^ testing is to obtain the required

plot of J-integral (J ) versus crack extension i^a) (Fig. 1) from the

measured quantities, which are the load (P ) on the specimen and the

relative displacement (6) of tvio measurement points located on the

specimen load line on opposite sides of the notch. The J value is

proportional to the area under the (p, 6) curve (Fig. 2), and can be

207
easily determined for any point along the curve. The shaded area in

Fig. 2 is proportional to J at the point labeled "a ." The principle

used to measure Aa is that the crack length, a, is related to the slope

of the specimen's (P, 6 ) curve for elastic conditions through the known

specimen compliance function [3]. The Aa measurement procedure is to

partially (about 10%) unload the specimen while recording load and

displacement changes (Figs . 2 and 3). Subsequent unloadings at increasing

crack extension values result in increased compliance. The crack extension

at a given unloading is the difference between the crack length measured

in that unloading minus the initial crack length, measured by an unloading

carried out before crack extension.

EQUIPMENT AND PROCEDURE

The loading apparatus used for computer-aided J-integral testing

was described previously by Fowlkes and Tobler [4]. The loading is controlled in conventional fashion by the operator. The computer is used

only in data acquisition and reduction. An analog ramp generator controlled

by the operator applies a smoothly varying signal to the test machine

control circuitry to allow unloading-reloading cycles as necessary. The

additional equipment needed for computer-aided testing includes a multichannel analog-digital conversion system, the computer itself, a CRT

terminal, a dual floppy disk data storage unit, and a digital plotter

(Fig. 4).

Two signals are used to direct the operation of the data acquisition

system. One is derived from the output of the ramp generator by which

the operator initiates the unloading-reloading cycle. When the ramp

generator is activated, a signal transmitted to the processor commands

it to begin to acquire and store load and displacement change data

during the unloading-reloading cycle. Data are stored until the load

again reaches the value it had before the unloading. The test continues

208
through as many unloading cycles as desired until the operator activates

the quit signal, which directs the processor to await the operator's

command (entered at the CRT terminal) to stop the data acquisition

procedure and terminate the test.

The computer-aided J-integral test facility at NBS is similar to that

described by Joyce and Gudas [3]. The differences are in the specific

types of equipment used and in the real-time outputs. "Real time" is used

here to mean "during the course of the test," as opposed to after the test

is completed.

The procedure for extracting J and Aa from the directly measured

load and displacement on a deeply notched compact tensile specimen is as

fol 1 ows

The electrical signals from the load cell and displacement gage are

amplified by the electronics supplied with the closed-loop servo-hydraul ic

apparatus used for the test. A plot of load against displacement similar

to Fig. 2 is produced by an x-y plotter throughout the course of the

test. Either signal may be displayed on a digital readout. The plot

and readout are used by the operator to monitor the progress of the

test.

The amplified load and displacement signals are also introduced

into an analog-to-digital conversion system to permit calculation of J

and Aa. At this point, the overall levels of the signals range between

0 and 10 V. These signals are converted to digital form with a resolution

of 16 bits; the least significant bit represents about 0.3mV. The

overall load and displacement values are used for calculating the instantaneous

J value, which is proportional to the area under the load-displacement

_3 curve. Stored overal 1 -displ acement points are separated by 2.5 x 10 mm.

Smaller observed overall-displacement increments are not stored. Computations of the instantaneous J value are made v/hen the overall displacement

209
has increased sufficiently from the previously stored value.

During the incremental unloading required to measure Aa, displacement

changes of the order of 0.05 mm and load changes of the order of 5 kN

must be correlated accurately. These changes correspond to a change in

the overall displacement signal of about 70 mV and a change in the load

signal of about 1 V. Because the analog-to-digital converter used has

only 16-bit resolution, the displacement signal cannot be converted

directly to digital form with precision sufficient for accurate measurement

of this displacement change. The measurement of the small displacement

changes which occur during the unloading-reloading cycles is facilitated

by regarding the overall displacement as consisting of small changing

displacements superimposed on a large reference displacement value. The

measurement is accomplished by cancelling out the reference part, then

converting the small displacement changes to digital form with high

resolution. To achieve this cancellation, the overall displacement is

converted to digital form at the beginning of each unloading-reloading

cycle. This initial value is treated as the reference part.

To cancel out this reference part of the displacement signal, a digitalto-analog converter with 14-bit resolution is used to apply a voltage

equivalent to the reference part of the displacement signal to one

terminal of a differential high-sensitivity analog-to-digital conversion

channel. The overall displacement signal is applied to the other input

terminal of this channel. This results in the analog-to-digital conversion

of a voltage equal to the difference between the instantaneous value of

the displacement signal and its value at the beginning of the unloading.

This difference signal is the required change-of-displacement signal

and is converted to digital form with 16 bit resolution of its 80 mV

range. The least significant bit represents about 2.5 yV, which is

quite sufficient. The load change signal is produced in a similar

210
i.ianner using another digital -to-analog converted signal and converted

to digital form with 16-bit resolution at the 1,25-V level. Here the

least-significant bit represents about 40 yV, which is again quite

sufficient.

Approximately 80 load-change and displacement-change data points

are obtained for each unloading-reloading cycle over a 30 s period.

These data are plotted during the unloading-reloading cycle using a

digital plotter, forming a plx)t like Fig. 3. From this plot, the presence

of excessive noise or other improper behavior in the signals can be

easily observed. The load-change and displacement-change data are

correlated using a least-squares fit. From the resulting slope, the

instantaneous crack length is calculated using the known expression for

the compliance of the specimen.

Each (J, Aa) datum is plotted by the digital plotter as soon as it

is available, forming a plot similar to Fig. 1. This plot allows convenient monitoring of the progress and quality of the test in real time.

The values are also displayed on a CRT terminal for transcription. This

step will be replaced by an automatic print-out when facilities become

available. The 1 east-squares-fit and data output procedure takes about

twelve seconds.

The record of the completed test consists of the (J, Aa), (P, 6),

and (aP, a 6) plots, the transcribed (J, Aa) data, and three files stored

on a floppy disk during the test, one containing the (J, Aa) values and

a few other key data, one containing all the (P, s) data, and one

containing all the (aP,A 6) data.

CONCLUDING COMMENTS

This system has been in use for six months and has proved quite

satisfactory. In comparison with the mul ti pi e- specimen and noncomputeraided single-specimen techniques previously used, it has eliminated

211
about 4 h of data reduction per test, and has provided a real-time

readout of the test progress and quality. The cost of the complete

data-acquisi tion system, including analog-to-digital converters, processor, digital plotter, and CRT terminal, was about $25,000.

REFERENCES

1. J. D. Landes and J. A. Begley, "Test Results form J-Integral

Studies: An Attempt to Establish a Testing Procedure," in

Fracture Analysi

, ASTM STP 560, American Society for Testing and

Materials, Philadelphia, PA (1974), pp. 170-186.

2. G. A. Clarke, W. R. Andrews, P. C. Paris, and D. W. Schmidt,

"Single Specimen Tests for Determination," in Mechanics of

Crack Growth , ASTM STP 590, American Society for Testing and

Materials, Philadelphia, PA (1976), pp. 27-42.

3. J. A. Joyce and J. P. Gudas, "Computer Interactive Jj^ Testing of

Navy Alloys," in Elastic-Plastic Fracture , ASTM STP 590, American

Society for Testing and Materials, Philadelphia, PA (1979), pp.

451-468.

4. C. W. Fowlkes and R. L. Tobler, "Fracture Testing and Results for a

Ti-6A1-4V Alloy at Liquid Helium Temperature," Eng. Fract. Mech. £

(1976), pp. 487-500.

212
009

(uj/N>i) nvao3iNi-r

<T3 O

CL

>1 o

+-> 4->

D. 4-J C7) (O -I— c -a u_

CD

213

CRACK

EXTENSION,

Aa

(mm)
214

DISPLACEMENT

CHANGE,

A8

(mm
215

Fig.

3.

Load

Change

plotted

against

displacement

change

for

two

selected

unload-reload

cycles.
0

0

O

Q. E

I - O (0 O

Oj C CO

cr.

216

Computer-aided

J-inteoral

test

apparatus.
TENSILE AND NOTCH-TENSILE PROPERTIES

AT ROOM TEMPERATURE AND 4K

OF WELDS IN ALUMINUM ALLOY 5083-0 PLATE

ALCOA

217
i
TENSILE AND NOTCH-TENSILE PROPERTIES

AT ROOM TEMPERATURE AND 4K

OF WELDS IN ALUMINUM ALLOY 5083-0 PLATE

R. A. Kelsey

L. N. Mueller

ENGINEERING PROPERTIES AND DESIGN DIVISION

ALCOA LABORATORIES

ALCOA CENTER, PA 15069

219
ABSTRACT

Tensile and notch- tens i 1 e data at room temperature and were obtained for

butt welds in 51 nn thick 5083"0 plates donated by five companies. The welds

were made by the conventional gas-metal arc welding (GMAW) process and by a high desposition rate GMAW process.

The investigation shows that the data are fairly consistent for welds made by

the various companies, despite the fact that there were considerable variations in welding procedures.

The strengths of the the welds are appreciably higher at ^ K than at room temperature. Notch yield ratios, which are an indication of fracture toughness, are quite high, indicating that the welds are very tough.

220
Ahuminum alloy base metals and weldments have excellent properties at lov% temperature and therefore have been widely used in cryogenic structures [ 1 ]

. The principal application in recent years has been in the use of alloy 5083 -O

in the construction of welded 35nn diameter spheres for the shipboard transportation of liquefied natural gas [2-5]. These spheres are made by joining

together segments in plate thicknesses ranging from 51 to 203mm. As a result

of these applications considerable tensile, fracture toughness, and fatigue

data have been obtained on base metal and welds in 5083 -O plate at temperatures down to 77 K [6-25]. However, relatively little data has been obtained

on 5083 base metal at kK [7, 11, 23] and the only data available on the

properties of welds in 5083 at kK is some work done at Alcoa Laboratories

about ten years ago [12].

The program reported herein was sponsored by National Bureau of Standards

and carried out at Alcoa Laboratories to obtain additional data on tensile r,nd

toughness properties of welds in 5083~0 plate at 4K.

The program had the following objectives:

1. Develop tensile and notch- tens i 1 e data on welds in 5083"0 at A K and

compare with room temperature data;

2. Compare the tensile and notch- tens 1 1 e properties of welds made by

different companies;

3 . Compare the tensile and notch- tens i 1 e properties of welds made by

the conventional gas-metal arc welding (GMAW) process with those

made by the high deposition rate (HDR) GMAW process.

This report also Summarizes some previously published data on tensile

properties of base metal and welds in aluminum alloys 2219

and 7039 at 4 K [11,12, 26]. These alloys, in addition to having adequate

toughness, have significantly higher yield snd tensile strength than 5083.

II. MATERIALS

Welded panels of aluminum alloy 5 O 83 -O plate for this program were donated

by the following companies:

Alcoa Laboratories, Aluminum Company of America

Chicago Bridge & Iron Company

Kaiser Aluminum and Chemical Corporation

P i ttsburgh-Des Moines Steel Company

Reynolds Metal Company

Each company was asked to donate two panels in 51 mm thick 5083"0 plate welded

by the GMAW process using either 5556 or 5183 electrode, with each panel

welded by a different welder. Alcoa and Kaiser also each donated two panels

welded by a high deposition rate GMAW process using 5183 electrode, again

using a different welder for each panel.

Numbers in brackets refer to references listed at end of text.

221
Figure 1 shows that the room temperature tensile properties of the alloy

5083“0 plates furnished by the five suppliers exceeded the guaranteed minimum

va 1 ues

III. WELDING PROCEDURES

The selection of welding procedures was left up to the various companies.

They were merely requested to weld 51mni thick 5083"0 plate in the flat

position by the GMAW process using a double-vee butt joint and either 5183 or 5556 electrode. It was specified that all panels should meet ASME radiographic

qual i ty standards.

Table I summarizes! the welding procedures used by the various companies. These

data show that there was considerable variation among the six suppliers. The

number of passes used to complete the welds made by the conventional GMAW process

ranged from 16 to 55, two companies used 3/32-in. diameter electrode; the

others used l/l6-ln. diameter. Voltages ranged from 24 to 34 volts and current

ranged from 230 to 390 amps. One company used pure argon shielding gas, the

others used a mixture of argon and helium, in various proportrons. Interpass

temperatures ranged from 150°F to 500°F. Welds were made automatically or semi-automatical ly; none were made manually. In the case of the high deposition rate welds, one company made the weld in four passes, using 1/8-in. diameter

electrode, 58O amps current and 34 volts; the other company used two passes,

3/l6-in, diameter electrode, 740 amps current, and 33 volts.

IV. TEST PROCEDURES

The tests at 4K were made in liquid helium in a cryostat described previously [ll].

Two types of tensile specimens were used to measure the cross-weld tensile

properties: room temperature and 4K tests were made using 12mm diameter

specimens while room temperature tests were also made using full-section

retangular specimens (Fig. 2). The 12 mm diameter specimens were located

at the mid-depth of the weld and had a 102 mm long uniform diameter test

section, which encompassed the root of the weld and the adjacent heataffected zone in the base metal on either side of the weld. Thus, fracture

occurred in the weakest region. The full-section tensile specimens were 25 mm wide and were tested with the weld reinforcement Intact. The weld joint

was located at the mid-length of the 279 mm long test section. Yield Strength

and elongation values at room temperature were based on deformation measurements made over a 51 mm gage length for the 12 mm diameter specimen and over

a 254 mm gage length for the full-section specimen. In the tests at 4K yield

strength was determined from load-head motion curves.

Cross-weld notch- tens i on specimens (Fig. 3) were tested at room temperature and

4K. The notch-tens ion test has been used for a number of years to evaluate

the relative notch toughness of various metals [27,28], including aluminum

alloys [ 29 , 30 ], but there are size limitations and conp 1 i cat i ons with such

things as fatigue precracking that make it less attractive than other

available procedures for measuring fracture toughness (K|q)[ 31]- However, it

has been demonstrated that there is reasonably good correlation between the

notch-yield ratio (NYR)* and K to warrant the acceptance of the notch-tension

test as a quality control screening of fracture toughness [ 32 ].

HtNYR = Notch- tens i 1 e strength

Tensile yield strength

222
The ASTM Tentative Method for Sharp-Notch Tension Testing of Cylindrical

Specimens (E602-76T) specifies two specimen diameters, namely 12 and 27 mm.

In extremely tough materials, such as base metal and welds in alloy 5083 -O,

the NYR relationship may suffer loss of sensitivity to because of general

yielding (on the cross-section containing the notch) prior to fracture. For

this reason, the larger diameter specimen (27mm) was used in this investigation

to reduce net section yielding and thus make the test results more discriminating, in terms of toughness indices.

V. RESULTS AND DISCUSSION

The main purpose of this investigation was to evaluate the effect of supplier

(welder), welding process, and test temperature on tensile and notch- tens i 1

properties obtained from specimens taken across the root of the weld. Obviously,

data obtained In this manner can only be considered to provide a relative

measure of weld joint performance since the results may be affected by metallurgical structure and physical or metallurgical discontinuities unique to the

root region. Therefore, to obtain a more representative evaluation of overall

(or average) joint performance, room temperature tensile tests were made of

full-section specimens (Fig. 2). Analysis of tensile properties obtained from

the full-section specimens (Table 2 and Fig. k) shows that at the 95 percent

confidence level there is no significant difference in the room tempera-ture

tensile properties between welds made by the five suppliers or between conventional GMA welds and high deposition rate GMA welds made by Suppliers 4 and 5-

The yield strengths range from 152 to 176 MPa, with a mean value of l6l MPa.

Tensile strengths range from 273 to 355 MPa, with a mean value of 314 MPa.

The minimum tensile strength specified by the ASME Boiler and Pressure Vessel

Code for butt welds in 51 nim 5 O 83 -O plate is 269 MPa. Elongations based on

a 10" in. gage length range from 7.7 to 18. 6 percent, with a mean value of 12.9

percent

CONVENTIONAL GMA WELDS

The results of tests of the tensile and notch- tens i 1 e specimens taken form

the mid-depth of the weld were evaluated by analysis of variance at the 95

percent confidence level.

Important statistically significant trends of data for the conventional GMA

welds. Table 3» Figs. 5"8, may be summarized as follows:

1. Supplier has a slight effect and temperature a relatively greater effect

on yield strength. The mean yield strength at 4K (201 MPa) is 23

percent higher than that at room temperature.

2. Supplier has a slight effect and temperature a relatively greater effect

tensile strength. The mean tensile strength at 4K (412 MPa) is 44 percent

higher than that at room temperature. The variability of the data is

greater at 4K than at room temperature.

3 . Supplier and temperature a slight effect on NYR. The mean value at 4K

( 1 . 80 ) is 8 percent lower than that at room temperature. Note that

although both yield strength and notch- tens i 1 e strength increase with

decreasing temperature, the percent increase in yield strength is greater

and therefore the net effect is that NYR decreases with decreasing temperature. The NYR values obtained for these welds are indicative of a very

tough material. By way of comparason, high toughness aluminum alloys

used in aircraft structures have NYR's of I .5 or less.

2i3
HIGH DEPOSITION RATE MS. CONVENTIONAL WELDS

Important statistically significant trends based on a comparison of conventional

GMA welds and high deposition rate GMA welds from Suppliers A and 5. Table A and Figs. 9“12, may be summarized as follows:

1. Supplier has a slight effect and temperature a relatively greater effect

on yield strength of HDR welds. At room ten^erature, the HDR welds have

lower yield strength but at AK the yief^are s imi lar^^^lds made by both

processes. It is important to remember that the tests of full-section

specimens show no significant difference in room temperature yield strength

of welds made by the two processes.

2. Supplier has a slight effect and temperature a relatively greater effect

on tensile strength of HDR welds. Tensile strengths of conventional and

HDR welds are similar at a given temperature. The variability in tensile

strength of HDR welds is g,reater at AK.

3. NYR's are slightly higher at room temperature and slightly lower at AK

for the HDR welds, as compared to conventional welds, but all values are

still indicative of a very tough material.

High deposition rate welding processes have several obvious aclwantages. Time

for welding thick plate can be reduced 50 to 80 percent compared to conventional

GMA welding time. Gas consumption is significantly reduced as compared to

conventional welds. Greater joint gaps can be tolerated without necessity of a backup. High deposition rate welds can be made in the horizontal and vertical

positions, as well as in the flat position.

COMPARISON OF BASE METAL AND WELDS IN ALLOYS 7039,2219 AND 5083

Tensile properties at room temperature and AK for base metal and conventional

GMA welds in nominal 51 nim thick plate in alloys 7039“T63, 2219"T851 and

5083"0 are shown in Table 5. These data show that both 2219 and 7039 base

metals have significantly higher yield and tensile strengths than does 5083 at

AK. Welds in 2219 also have significantly higher yield strength than 5083. No

AK data are available for welds in 7039 or any of the other Al-Zn-Mg alloys.

NYR data available for base metal and welds in 2219 and 7039 indicate that

these materials have adequate toughness at 4 K.

Al-Zn-Mg alloys are of particular interest because their relative insensivity

to cooling rate following quenching from solution heat treating temperature

results in high tensile properties in thick sections. This, and their ability

to naturally age harden, also results in good weld strengths. Some work has

been done in developing filler alloys for welding Al-Zn-Mg alloys which shows

that it is possible to obtain weld strengths higher than those obtained using

commercially available filler alloys such as 5356. Work is needed to develop

these filler metals to provide optimum strength at AK, as well as to develop

optimum procedures for welding thick sections in these materials. Furthermore,

it may also be possible to develop Al-Zn-Mg base metals with improved

strengths at AK.

224
VI-. SUMMARY AND CONCLUSIONS

This report contains tensile and notch- tens i 1 e data at room temperature and

for butt welds in 51 mm thick aluminum alloy 5083-O plates donated by five

companies. Welds were made the conventional gas-metal arc welding (GMAW)

process and by a high desposition rate GMAW process.

The results of this investigation, which were eval uated stat i st i cal 1 y by the

method of variance, show that:

1. Properties of welds made by the five companies were quite consistent,

despite the fact that there. were considerable variations in welding

procedures

2. At 4K the mean yield strength (201MPa) and the mean tensile strength

(421 MPa) are 23 and 44 percent higher, respectively, than those at room temperature.

3. Notch yield ratios, which are an indication of fracture toughness , are quite high, indicating that the welds are very- tough.

4. Properties of welds made by the high deposition rate process, which can reduce welding time for thick plates by 50 to 80 percent, are comparable to

properties of welds made by the conventional GMAW process.

225
REFERENCES

1. Rosenthal, Alan, "Aluminum Welding Advances Launch New Wave of Commercial,

Naval Craft," Modern Metals, September, 1967.

2. "Giant Aluminum LNG Sphere 12 Stories High Arrives at General Dynamics Quincy

Yard," Maritime Reporter and Engineering News, January 15, 1977, pp. 8-9.

3. "Aluminum in Hot Pursuit of Cryogenics," Welding Design and Fabrication, 1974.

4. "LNG Keeps Growing Around the World," Pipeline and Gas Journal, June, 1976,

pp. 18-22.

5. Meyer, H. and Paetow, K. H., "125,000 m^ LNG/LPG Tanker at HDW," Schiff and

Hafen, 11, 1975, pp. 989-997.

6. Kaufman, J. G., and Nelson, F. G., "Cryogenic Temperatures of Fatigue Strength

of Al-Mg Alloys," Space/Aeronautics, July 1962.

7. Rice, L. P., Campbell, J. E., and Simmons, W. F., "The Tensile Property Evaluation

of One 5000-Series Aluminum Alloy at the Temperature of Liquid Helium," Advances

in Cryogenic Engineering, Vol. 8, 1963, pp. 671-677.

8. Kaufman, J. G., and Holt, M., "Fracture Characteristics of Aluminum Alloys,"

Aluminum Company of America, Alcoa Laboratories, Technical Paper No. 18, 1965.

9. Kaufman, J. G., and Wanderer, E. T., "Aluminum for Cryogenic Applications," Machine

Design, November 11, 1965.

10. Nelson, F. G., Kaufman, J. G., and Holt, M., "Fracture Characteristics of Welds

in Aluminum Alloys," Welding Journal, July 1966.

11. "Kaufman, J. G., Bogardus, K. O., and Wanderer, E. T., "Tensile Properties and

Notch Toughness of Aluminum Alloys at -452°F in Liquid Helium," Advances in

Cryogenic Engineering, Vol. 13, 1968, pp. 294-308.

12. Nelson, F. G., Kaufman, J. G., and Wanderer, E, T., "Tensile Properties and Notch

Toughness of Groove Welds in Wrought and Cast Aluminum Alloys at Cryogenic

Temperatures," Advances in Cryogenic Engineering, Vol. 14, 1969, pp. 71-82.

13. Nelson, F. G., Kaufman, J. G., and Wanderer, E. T., "Tear Tests of 5083 Plate and

of 5183 Welds in 5083 Plate and Extrusions," Vol. 15, Advance in Cryogenic Engineering, 1970.

14. Kaufman, J. G., Nelson, F. G., and Wanderer, E. T., "Mechanical Properties and

Fracture Characteristics of 5083-0 Products and 5183 Welds in 5083 Products,"

Vol. I, Proceedings of the XIII International Congress of Refrigeration, Washington,

DC, 1971.

Holt, M., and Kaufman, J. G., "Indices of Fracture Characteristics of Aluminum

Alloys Under Different Types of Loading," Current Engineering Practice, Vol. 16,

No. 3, 1973.

226

15.
16. Nelson, F. G., and Kaufman, 3. G., "Fracture Toughness of Plain and Welded 3-in.

Thick Aluminum Alloy Plate," ASTM STP 536, Progress in Flaw Growth and Fracture

Toughness Testing, 1973, pp. 350-376.

17. Saunders, W. W., 3r., and Gannon, S. M., "Fatigue Behavior of Aluminum Alloy

5083 Butt Welds," Welding Research Council Bulletin 199, October 1974.

18. Kelsey, R. A., Nordmark, G. E., and Clark, 3. W., "Fatigue Crack Growth in Aluminum

Alloy 5083-0 Thick Plate and Welds for Liquefied Natural Gas Tanks," ASTM STP

556, Fatigue and Fracture Toughness - Cryogenic Behavior, 1974, pp. 159-185.

19. Kaufman, 3. G., Nelson, F. G., and Wygonik, R. H., "Large-Scale Fracture Toughness

Tests of Thick 5083-0 Plate and 5183 Welded Panels at Room Temperature, -260

and -320°F, ASTM STP 556, Fracture and Fracture Toughness - Cryogenic Behavior,

1974, pp. 125-158.

20. Kelsey, R. A., Wygonik, R. H., and Tenge, P., "Crack Growth and Fracture of Thick

5083-0 Plate Under Liquefied Natural Gas Ship Spectrum Loading," ASTM STP

579, Properties of Materials for Liquefied Natural Gas Tankage, 1975, pp. 44-79.

21. Kaufman, 3. G., and Kelsey, R. A., "Fracture Toughness and Fatigue Properties

of 5083-0 Plate and 5183 Welds for Liquefied Natural Gas Applications," ASTM

STP 579, Properties of Materials for Liquefied Natural Gas Tankage, 1975, pp.

138-158.

22. Bogardus, K. O., and Malcolm, R. C., "Design Stresses for Aluminum Alloy 5083-0

and 5183 Welds at Cryogenic Temperatures," Ibid, pp. 190-204.

23. Tobler, R. L., and Reed, R. P., "Fracture Mechanics Parameters for a 5083-0

Aluminum Alloy at Low Temperatures," 3. of Engr. Materials and Technology, Vol.

99, Oct. 1977, pp. 306-312.

24. Kelsey, R. A., "Fatigue and Fracture Characteristics of Al-Mg Butt Welds," Aluminium,

October 1977, pp. 600-606.

25. Kelsey, R. A., "Mechanical Properties of USA/USSR Al-Mg Plate and Welds for

LNG Applications," Vol. 24, Advances in Cryogenic Engineering, 1978.

26. Kaufman, 3. G. and Wanderer, E. T., "Tensile Properties and Notch Toughness of

Some 7XXX Alloys at -452®F," Advances in Cryogenic Engineering, Vol. 16, 1971,

pp. 27-36.

27. Sachs, G., Ebert, L. 3., and Brown, W. F., 3r., Transactions, American Institute

of Mining, Metallurgical, and Petroleum Engineers, Iron and Steel Divisions, 1948,

p. 424.

28. Baker, 3. F. and Tipper, C. F., Welding Research Council Bulletin, No. 37, 3une

1957.

29. Kaufman, 3. G. and 3ohnson, E. W., Advances in Cryogenic Engineering, Vol. 6,

1960, pp. 637-649.

30. Kaufman, 3. G. and Holt, Marshall, "Fracture Characteristics of Aluminum Alloys,"

Technical Paper No. 18, Alcoa Laboratories, 1965.

227
31. Brown, W. F., Jr., "The Circumferentially Notched Cylindrical Bar as a Fracture

Toughness Test Specimen," presented at the G. R. Irwin Symposium on Linear

Fracture Mechanics, Lehigh University, 29-30 July 1974.

32. Kaufman, J. G., Sha, G. T., Kohm, R. F., and Bucci, R. J., "Notch-Yield Ratio as a Quality Control Index for Plane Strain Fracture Toughness," ASTM STP 601, 1976,

pp. 169-190.

228
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244 FU^.I2
NOTE ON A SIMPLE TEST METHOD FOR

ESTIMATING J.

Ic National Bureau of Standards

245

Note on A Simple Test Method for

Estimating J.

Ic

T. A. Whipple and H. I. McHenry

National Bureau of Standards

Boulder, Colorado

ABSTRACT

Fracture toughness testing is generally a time-consuming and expensive

procedure; therefore, there has been a significant amount of effort

directed toward developing an- inexpensive and rapid method of estimating

the fracture toughness of materials. In this paper, a simple method

for estimating through the use of small, notched, bend bars is

evaluated. The test only involves the measurement of the energy

necessary to fracture the sample. Initial tests on Fe-18Cr-3-Ni-l 3Mn

and 304LN stainless steel at 76 and 4 K have yielded results consistent

with other fracture toughness tests, for materials in the low- to mediumtoughness range.

247
INTRODUCTION

The fracture toughness of stainless steel welds at 4 K varies

significantly depending on the weld metal chemistry and the welding

process and procedures, as shown by Whipple, McHenry, and Read in this

report (p.l67). Consequently, it is necessary to qualify welding procedures

and consumables for liquid helium service. Weld qualification testing

is currently done for the various magnet programs by NBS using the

procedures described by Read in this report (p. 205). Read's method is

suitable for research purposes, but it is too expensive and complex for

routine qualification testing. A simple and inexpensive method of

measuring or estimating at 4 K is needed to permit fabricators and

independent laboratories to conduct qualification tests.

The idea of simply conducting a Charpy test at 4 K -- or better yet

at 77 K -- is not regarded as suitable. In a Charpy impact test at

4 K, the rapid deformation prior to fracture causes adiabatic heating of

the specimen and invalidates the results. This problem is inherent to

all structural alloys because of their low specific heat at 4 K.

Previous work at NBS [1] indicates that testing at higher temperatures,

e.g., in liquid nitrogen at 77 K, does not provide a reliable indicator

of toughness at 4 K.

Begley [2] developed a method of estimating toughness from the

fracture energy of notched bend bars based on the following relationship:

^T = Jt + a(W-a), for p

= constant

BlTr^ W-a

where Ej = fracture energy

Jlc = fracture toughness

248
p

= notch root radius

B = thickness

W = width

a = crack length

a

= empirical constant

This relationship was derived on the assumption that fracture occurs by

a ductile tearing mode after general yielding. These assumptions are

reasonable for austenitic stainless steel weldments tested using deeply

notched bend bars. In this note a limited experimental program to

evaluate Begley's method is discussed.

Material

The test materials were 304LN stainless steel and Fe-18Cr-3Ni-l 3Mn

for which the (J) was known from previous tests [3]. These materials

were selected because of the wide range of fracture toughness values,

measured by either the single- or multiple-specimen test methods, as

shown in Table I

Experimental Procedure

The samples used were 10-mm x 10-mm (Charpy-impact-size) bend bars

with two different notches. One notch was 5-mm (0.197-in) deep with a

root radius of 0.5-mm (0.02-in) the other was 7. 5-mm (0.296-in) deep

with a root radius of 0.25-mm (0.01-in). This gave both samples a (Wa)/p of 10. Impact tests were conducted at 76 K and slow-bend tests

at 76 and 4 K. The specimens were oriented such that the notch was in

the rolling direction perpendicular to the plate surface.

Impact Tests : Impact tests were conducted according to the procedures

described in ASTM E23, with the exception of the notch geometry as

described above.

249
S1ow-Bend Tests: The fixture for slow-bend testing was designed such

that it fit into the cryostat used for tensile testing, which has been

described by Reed [4]. The loads are applied to the sample through

three 6.35-mm (0.25-in) diameter pins.

The load versus time curve is recorded on a strip chart recorder.

The time scale on the test record is converted to deflection by using

the known machine speed. The total energy for the test is equal to the

area under the load deflection curve. To obtain the energy used to

fracture the sample, the energy stored in the machine at the fracture

load is subtracted from the total energy. The value for the machine

energy as a function of load was found by bending an unnotched sample in

the test fixture at the test temperature. The deflection of the unnotched sample (calculated from beam equations) was subtracted from the

total deflection to obtain the machine deflection as a function of load;

which is then converted to energy.

Results and Discussion

Table I gives the results obtained from the impact and slow bend

tests, along with the J. results from the standard tests. The Jt

results for Fe-18Cr-3Ni-13Mn were obtained by Read and Reed [3] on the same

heat of material by the multiple-specimen J-integral test. The

results for 304LN were obtained on the same heat of material by the

single-specimen J-integral test.

Figure 1 is the plot of the normalized fracture energy versus the

remaining ligament showing the extrapolation to zero remaining ligament

which yields the estimate of Jj^. These values are given in Table I.

250
It is interesting to note that the normalized fracture energies

for impact tests are significantly higher than those for the correspondi ng

slow-bend tests. However, the extrapolation yields values that are

very similar. This is probably due to vibrational and other losses in

the impacting process, which are a function of the energy required to

break the sample and are negated by extrapolation to zero ligament.

Figure 2 is a comparison of the values measured by the singleor multiple-specimen tests and. the estimated values obtained in this

study. It shows that the estimation is excellent at low toughness

values and becomes conservative as the toughness increases. This is not

surprising, considering the size of sample used for these tests. According

to the proposed validity criteria for testing of;

^Ic

B, (W-a) > 25 —

The maximum these tests should be 0.0002 mm (0.0079 in).

This means that at 4 K for 304LN the maximum is 154 kJ/m^ (900 in*lb/in^)

and for Fe-1 8Cr-3Ni-l 3Mn it is 229 kJ/m^ (1310 in*lb/in^) at 76 K. By

this criterion, all the 304LN tests are invalid and all the Fe-1 8Cr-3Ni-l 3Mn

tests are valid.

From the limited data obtained to date, Begley's test method appears

to be a promising estimation method. When Charpy-sized specimens

are used, it is particularly suitable for materials in the low- to

medium-toughness range. Thus, it may prove very useful for testing

2 stainless steel weldments, which generally have a below 175 kJ/m

(1000 in*lb/in“) at 4 K. The application of this test to stainless

steel weldments is currently under study.

251
References

1. D. T. Read, H. I. McHenry, P. A. Steinmeyer, and R. D. Thomas, Jr.,

"Metallurgical Factors Affecting the Toughness of 316L SMA Weldments

at Cryogenic Temperatures ," in Material Studies for Magnetic Fusion

Energy Applications at Low Temperatures-I

, NBSIR 79-1609, 313 to

352, (1979).

2. J. A. Begley, "An Estimate of Ductile Tearing Jj^ Toughness from

the Fracture Energy of Notched Bend Bars," unpublished.

3. D. T. Read and R. P. Reed, "Toughness, Fatigue Crack Growth, and

Tensile Properties of Three Ni trogen-Strengthened Stainless Steels

at Cryogenic Temperatures," in Material Studies for Magnetic Fusion

Energy Applications at Low Temperatures-I , NBSIR 78-884, 93 to 153,

(1978).

4. R. P. Reed, "A Cryostat for Tensile Tests in the Temperature Range

300 to 4 K," in Advances in Cryogenic Engineering, Vol . 7 , Plenum

Press, New York, 1962, pp. 448-454.

252
-a ^

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1

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253

*Measured

by

singleor

multiple-specimen

J-inteqral

test.
LIST OF FIGURES

Fi gure 1 .

Figure 2,

Normalized fracture energy as a function of the ligament

width with the extrapolation to zero remaining ligament.

Comparison of the J. values obtained by the estimation

procedure and by the^single or multiple specimen tests.

254
Figure 1. Normalized fracture energy as a function of the ligament

width with the extrapolation to zero remaining ligament.

255
MEASURED

(kj/m

500

400

C\J 300

200

1

100

304LN ^ 76 K (Impact)

O 76 K (Slow Bend)

4 K (Slow Bend)

Fe-18Cr-3Ni-13Mn

A 76 K (Impact) • 76 K (Slow Bend)

4 K (Slow Bend)

/

/

/

/

/

/

/

/

/

1/

100 200 300

Jj^ESTIMATED

400 500

Figure 2. Comparison of the Jjq values obtained by the estimation procedure

and by the single or multiple specimen tests.

256
Nonmetallics

257

NONMETALLICS FOR MAGNET SYSTEMS

NationaT Bureau of Standards

259

NONMETALLICS FOR MAGNET SYSTEMS

R. E. Schramm, M. B. Kasen, H. M. Ledbetter, J, G. Must, and A. F. Clark

Fracture and Deformation Division

and

Thermophysical Properties Division

National Bureau of Standards

Boulder, CO 80303

INTRODUCTION

NBS has undertaken the task of characterizing the mechanical and

physical parameters of G-IOCR and G-llCR glass cloth/epoxy composites

between room temperature and 4 K. These materials meet current NEMA

G-10 and G-11, MIL-P-181 77C, and Federal LP-509 specifications. In

contrast to proprietary NEMA grades, the CR laminates are produced to

common specifications covering all aspects of components and procedures.

These products have minimum cryogenic property variability and are

commercially available from several manufacturers

Determination of mechanical properties involved: 1) tensile tests

to measure Young's modulus, Poisson's ratio, ultimate strength, and

failure strain, 2) compression strength tests, and 3) interlaminar shear

strength tests. These were done for both major directions in the glass

cloth plane (warp and fill). Compression strength was also measured in

the direction normal to this plane. Ultrasonic resonance offers unique

possibilities to examine elastic properties and was used to measure

modulus and internal friction.

Thermal properties measured were conductivity and contraction.

Values of electrical resistivity and dielectric strength provided by a

subcontractor were included in the last report. C. J. Long of ORNL

conducted radiation resistance tests.

261

CRYOGENIC PROPERTIES OF FILAMENTARY-REINFORCED

COMPOSITES: AN UPDATE

National Bureau of Standards

2G3

CRYOGENIC PROPERTIES OF FILAMENTARY-REINFORCED COMPOSITES: AN UPDATE

M. B. Kasen

Fracture and Deformation Division

National Bureau of Standards

Boulder, Colorado 80303

ABSTRACT

Current knowledge of the effect of cryogenic temperatures on the

mechanical, elastic, thermal and electrical properties of reinforced

plastic laminates is reviewed. Directions of fundamental research and

of test method development are defined and discussed. Recommendations

are made for effective future research.

It is confirmed that conventional laminates developed for room

temperature service are viable engineering materials at cryogenic temperatures.

Averages of mechanical and elastic properties increase on cooling. But

increased variability in strength properties on cooling remains a problem.

This variability reflects the increasing notch sensitivity of the matrix

resins on cooling.

It is concluded that the greatest benefit would accrue from development

of matrix resin systems combining room temperature strength with improved

cryogenic toughness.

265
INTRODUCTION

The cryogenic properties of fiber-reinforced composites were last

reviewed by Kasen in 1975 (1-3). These reviews identified a number of

thermal and mechanical property trends for the room temperature to cryogenic temperature range for a variety of composite materials. A global

view of available data showed that cooling generally had a beneficial

effect on mechanical properties. No instance of serious property degradation

was reported. Cotton, glass, graphite, aramid and boron-reinforced

polymers and boron-reinforced aluminum therefore appeared to have potential

for cryogenic service.

But, it was equally evident that a primary obstacle to expanded use

of composites was the lack of a systematic data base relating cryogenic

properties to well characterized composite materials. Available data

often reported significantly differing temperature dependencies for

mechanical properties of a given type of composite material . This

reflected a diversity of data sources each having different program

objectives and using different experimental techniques. Property data

below liquid nitrogen temperature (76 K) were limited to 20 K. Almost

no data were available at liquid helium temperature (4 K). Some efforts

to improve cryogenic composite performance by flexibil izing the resin

system were reported. But results were inconclusive.

266
Recent years have seen substantial progress in experimental methodology.

Furthermore, developments in composite mechanical theory have defined

ways to produce a more systematic characterization of cryogenic composite

laminate properties. To a large extent, such refinements paralleled the

rapid development of overall composite technology.

Recent efforts to understand composite behavior better at cryogenic

temperatures have largely been stimulated by the need to develop power

generation systems as alternatives or improvements to conventional

systems (4,5). Current worldwide interest in energy generation by

magnetic fusion energy (MFE) and magnetohydrodynamic (MHD) technology

and in improving the efficiency of rotating machinery has required

construction of large superconducting magnets which must function for

many years under demanding service conditions. This has placed increasing

demands on the performance of composite materials that serve as electrical

insulators and primary supports for the magnet windings. It has also

stimulated interest in the potential of advanced, high-modulus composites

as contributors to lighter weight and increased efficiency of cryogenic

power systems.

This paper provides an interim report on the progress since 1975 in

understanding the performance of composite materials at cryogenic

temperatures. The current data base is discussed within two composite

materials categories: 1) woven-fabric and random-mat reinforced composites

and 2) uniaxial fiber reinforced composites. The first category encompasses

relatively inexpensive laminates, often produced in bulk form by industrial

laminating firms. The second category addresses the most advanced.

267
high-performance types of laminates. Succeeding sections review the

status of fundamental research and of test method development. Finally,

the directions of effective future research efforts are considered.

WOVEN-FABRIC AND RANDOM-MAT REINFORCED COMPOSITES

OVERVIEW

Polymer-matrix composites reinforced with woven-fabric or randommat fibers are the least expensive of the laminate systems when they are

produced by industrial high pressure laminating methods. The low cost,

reasonable strength properties and low thermal and electrical conductivities have made such materials very popular for cryogenic structures.

Glass-fabric reinforced epoxy composites continue to dominate cryogenic interest. Best overall performance is obtained from laminates

using conventional, ful ly-reacted epoxy resin systems. Laboratory work

and practical experience have provided evidence that highly flexibilized

epoxy systems do not improve the cryogenic performance of glass-fabric

reinforced laminates. Furthermore it is becoming clear that such resin

systems are a distinct detriment to room temperature performance. The

failure of highly flexibilized systems is largely due to their high

thermal contraction which creates a high level of residual stress in the

cooled laminate, negating the inherent residual strain capability of the

matrix at cryogenic temperatures.

A substantial effort has been directed toward reducing the cryogenic

property variability of NEMA G-10 and G-11 types of glass-fabric reinforced

hi gh-pressure industrial laminates that are produced in large quantities

by several U.S. laminating firms. Such products are fabricated to meet

normal temperature electrical insulation performance criteria established

by the National Electrical Manufacturer ' s Association (NEMA). Manufacturers

268
are given considerable leeway in selection of specific laminate components

and in manufacturing techniques. Proprietary NEMA-grade laminates

produced by the various firms therefore display substantial variability

in those mechanical and thermal properties of greatest interest to the

designers of large superconducting magnets and other sophisticated

cryogenic apparatus.

Recognizing this problem, a cooperative program was undertaken

between the National Bureau of Standards (NBS), the NEMA organization

and the U.S. laminating industry to establish cryogenic grades of NEMA

G-10 and G-11 glass-fabric reinforced epoxy products. The resulting

laminates are designated G-IOCR and G-llCR. These CR-grade products are

high-quality, state-of-the-art products, differing from proprietary

products in that they are produced to closely controlled, uniform specifications

covering both component materials and manufacturing techniques. The CR

products meet existing NEMA, MIL and Federal specifications for conventional

G-10 and G-11 products. G-IOCR is a heat-activated, amine-catalyzed,

bisphenol A solid-type epoxy resin laminate reinforced with continuousfilament E-glass having a silane finish. G-llCR is an aromatic-amine

hardened bisphenol A liquid-type epoxy resin laminate reinforced in the

same manner. Such laminates are currently in commercial production and

have been incorporated into several prototype superconducting magnets.

The mechanical, thermal and electrical properties of G-IOCR and

G-llCR have been characterized by NBS and other U.S. laboratories (6).

The effect of neutron and gamma irradiation at 4 K on the mechanical and

electrical performance of the materials have been studied by the Oak

Ridge National Laboratory (7-8).

269
Similar concern for the proper selection of laminate materials for

use in superconducting magnet technology has led to a substantial body

of cryogenic composite characterization work on a series of European industrial laminates by the European Organization for Nuclear Research

(CERN) in Geneva, Switzerland (9). Mechanical and thermal property data

were obtained from 275 to 4 K on E-glass fabric-reinforced bisphenol A

epoxy, epoxy-novol ac , polyimide, phenolic and silicone laminates plus a

pitch-base graphi te/bi sphenol A and an arami d/bisphenol A laminate. The

careful documentation of the studied materials and the sophistication of

data acquisition and reduction methods were outstanding in this program.

A substantial contribution to understanding the effect of cryogenic

temperatures on the performance of glass-mat reinforced epoxy laminates

has been made by a leading industrial glass research laboratory (10).

Development of a practical nonmetal lie containment system for liquefied

natural gas required a careful study of the effect of cryogenic temperatures

on the integrity of the laminate serving as the primary LNG barrier.

This program generated an extensive mechanical and thermal property data

base to 76 K for both discontinuous and continuous strand glass-mat/

epoxy laminates. Performance of the laminate systems under severely

constrained thermal shock was also thoroughly investigated. Mechanical

and thermal properties of these laminates have been summarized in handbook form (11) along with the properties of other nonmetallic laminates

of interest to LNG technology.

Studies on the cryogenic temperature performance of glass-reinforced

polyester laminates have been reported by the University of Wisconsin

(12-13). The work was motivated by the need to develop an inexpensive

270
laminate for the major structural supports and thermal standoffs in very

large superconducting energy storage systems.

DATA ANALYSIS

The most recent data on the temperature dependence of the significant mechanical, thermal, and electrical properties of fabric and matreinforced polymer composites are summarized on Figures 1-10 and on

Table I. Data obtained at temperatures defined by the boiling points of

cryogens (76 K and 4 K) are presented in bar graph form to indicate lack

of knowledge of the intervening temperature dependence. The cryogenic

temperature values are taken as the average of the data at 76 and 4 K

except where there is evidence of a statistically significant property

change from 76 K to 4 K. Data obtained over closely spaced temperature

intervals are presented as continuous curves.

Strength Properties

The maximum tensile strengths in fabric-reinforced laminates correspond

to the warp direction of the fabric where glass content is the highest.

The generally lower glass content in the fill direction results in lower

strength. The G-IOCR and G-llCR laminates have a 42:32 warp-to-fill

ratio typical of the style of glass fabric used in high-pressure industrial

laminates manufactured in the United States. Lowest strength is developed

in a 45° orientation with respect to the warp and fill.

Conversely, a laminate reinforced with continuous or discontinuous

randomly oriented glass mat has essentially isotropic properties in the

edgewise direction, i.e., in the plane normal to the direction of

pressing during laminate cure. But this isotropy is obtained at the

271
expense of reduced laminate strength and stiffness compared with the

maximum values obtainable from fabric-reinforced laminates.

The most recent data on the effect of cryogenic temperatures on the

tensile, compressive and interlaminar shear strengths of fabric and matreinforced polymer composites are summarized on Figures 1-3.

Figure 1 summarizes ultimate tensile strength data. Except as

noted, all data are for the warp direction of fabric reinforcement and

for the edgewise direction of mat reinforcement. In part, the data

compare reported strengths of a series of E-glass fabric-reinforced

bisphenol A epoxy laminates having varying glass contents. Conventional

bisphenol A epoxy resins (diglycidyl ether of bisphenol A) and epoxy

novolacs (glycidyl ethers of novolacs) are most widely used in laminate

fabrication (14). The composition of E glass is given on Table II.

These data compared to strengths reported for continuous and

chopped-mat E-glass reinforced bisphenol A epoxy products, for E-glass

fabric reinforced polyimide, and for polyester and melamine laminates.

Available data on the cryogenic performance of epoxy laminates reinforced

with aramid fiber fabric (Kevlar 49) and with a pitch-base graphite

fabric (HM-S) are also documented along with the small amount of data

available for cotton fabric reinforced phenolic (phenol formal dehyde)

1 ami nates

Highest cryogenic tensile strengths were reported for the NEMA

6-11 CR laminate containing 67-72 weight percent glass. G-IOCR, having

64-68 weight percent glass, developed slightly lower strength. CERN (9)

reports that a commercial laminate of European manufacture containing

74.8 weight percent glass (Permaglas TE630) developed properties similar

272
to the CR grades, while other commercial laminates having glass contents

of 62.5-66.1 weight percent (Vetronite 432.81, Vetronite 432.86, Durostone

EPF2) developed slightly lower cryogenic strengths (bar 4, Fig. 1).

These data illustrate the importance of specifying glass content or of

buying laminates to a specified range of glass content whenever maximum

laminate performance is desired.

NBS data (6) on Figure 2 indicate that the tensile strength in the

fill direction of G-IOCR and G-llCR are 35-70% of that in the warp

direction. A similar strength anisotropy can be expected for the

other fabri c-reinforced laminates depending on the warp/fill ratio of

the fabric used in their construction.

The tensile strengths developed by the mat-reinforced E-glass/

epoxy products (bars 5 and 6, Fig. 1) are substantially lower than those

developed for the warp direction of fabric-reinforced E-glass/epoxy

products. But the strength difference is less when compared to the

fill direction of the fabric products. It is expected that the strength

of the mat products would exceed that of the fabric products in a 45°

direction from warp and fill.

CERN (9) reports that a commercial polyimide-matrix laminate

(Silirite 64 400) developed 7-10% lower average ter'sile strength than

did bisphenol A epoxy laminates having similar type and content of glass

reinforcement (bar 7, Fig. 1).

The data presented for glass fabric-reinforced melamine (bar 8,

Fig. 1) are typical values reported by the manufacturer (11). This

type of laminate is commercially produced as NEMA G-5. Data presented

273
for glass fabric-reinforced polyester (bar 9, Fig. 1) reflect average

data summarized in the literature for a variety of such laminates (11).

Glass reinforced laminates all strengthen substantially on cooling.

These data suggest that the average tensile strength of epoxy-matrix

laminates will increase 45-55% over the room temperature values, whereas

the polyimide, melamine and polyester laminates increase about 35-40%.

Despite its fairly wide use, little is known of the effect of

cryogenic temperatures on the strength of cotton-reinforced phenolic

laminates.* The data appearing in Figure 1 reflect one manufacturer's

published data on its proprietary NEMA type CE product (11). Strengths

obtained from other manufacturers' products or from the NEMA type LE

cotton/phenolic product may differ substantially. But there is no doubt

that the strength of the cotton/phenolic product is inferior to that of

the glass-reinforced products.

The CERN data (9) on the pitch-base HM-S graphite fabric (Thornel

PVCB 45) and on aramid fabric (Kevlar 49) reinforced epoxy laminates

summarized on Figure 1 represent the first cryogenic data obtained on

these materials. The aramid laminate displays a significant average

strength reduction on cool i ng--possi bly a manifestation of internal

stresses created by the very large negative coefficient of thermal

expansion of this fiber type. The comparatively low strength of the

graphite fabric reinforced product reflects the low strength of the

inexpensive pitch-based fiber used in its construction. A comparable

*The term Micarta is often used when referring to cotton-reinforced phenolic

industrial laminates. However, Micarta is the trade name designating all

industrial laminates produced by the Micarta Div., Westinghouse Electric Corp.

274
laminate reinforced with graphite fabric using fiber derived from the

more expensive polyacryloni tril e (PAN) precursors should develop about

twice the tensile strength of the pitch-based product.

The compressive strength data presented on Figure 2 reflect both

the warp direction and normal direction where such data are available.

The direction designated "normal" is normal to the plane of reinforcement

for fabric-reinforced laminates and is parallel to the pressing direction

during laminate cure for mat- reinforced laminates.

These data suggest that cryogenic applications requiring compressive

strength in the normal direction may be well served by some of the less

expensive fabric-reinforced laminates such as the melamine-matrix NEMA

G-5 (Bar 4, Fig. 2) or by mat-rei nforced laminates (Bar 3, Fig. 2).

These laminates, the G-llCR grade of glass/epoxy and the NEMA CE cotton/

phenolic products develop higher average compression strength in the

normal direction than in the direction of reinforcement. The NEMA

G-IOCR product is the only laminate for which the average warp direction

compressive strength is reported to exceed that of the normal direction.

The data for G-llCR, G-IOCR and for chopped-mat glass/epoxy (Bars 1,

2, and 3 Fig. 2) represent systematic studies of well characterized

materials. Data for NEMA G-5 and CE (Bars 4 and 6, Fig. 2) are manufacturers' typical data. G1 ass/polyester data reflect the average of

literature values.

NBS data (6) indicate that testing in the fill direction of the

fabric reinforcement reduces the average ultimate compressive strength

by about 30% for the G-IOCR product and by about 24% for the G-llCR

product.

275
Interlaminar shear strength, the resistance to shear between the

layers of a layered laminate, has become an important design property in

large superconducting magnets where glass-fabric reinforced composites

are used for electrical insulation and for structural support of superconducting windings. The interlaminar shear strength of the selected

materials must be adequate to resist electromagnetic forces within the

magnet. The guillotine test method described by ASTM method D2733-70 is

the most appropriate. However, the application of this method as described

in ASTM standards has been found to yield extremely variable results.

In response to the needs of the MFE research efforts, NBS and the MIT

Francis Bitter National Magnet Laboratory have refined the test method

as described in a subsequent section of this paper.

The temperature/orientation dependence of interlaminar shear in a

G-IOCR glass-fabric reinforced epoxy industrial laminate appearing on

Figure 3 has been reported by MIT (15). NBS data for the warp (0°)

direction are in agreement with these data. NBS data for the warp

direction of G-llCR noted by circles on Fig. 3 indicate a somewhat lower

average interlaminar shear strength for this material.

Data on the average change in laminate strength properties on

cooling are useful in comparing cryogenic and room temperature performance.

But such data can be misleading from a design standpoint. Here the

relevant data are the lower bound of performance--not the average. It is

frequently observed that the variability in strength properties increases

with increases in average strength on cooling. The variability appears

to be less for the lower strength fabric and mat-reinforced laminates

than for the higher performance uniaxial laminates discussed in a later

276
section; nevertheless, it is a factor to consider in all laminate types.

Variability arises intrinsically due to embrittling and increased notch

sensitivity of the polymer matrix in a cooled laminate. It is therefore

a larger factor in properties such as shear and compression which are

most dependent on matrix properties. Variability is a strong function

of the quality of laminate manufacture. In particular, high void contents

and resin rich areas must be avoided. Least variability will be experienced

with laminates fabricated under the close control exercised by large

industrial laminating firms, particularly when a control 1 ed-specification

products such as the CR grades are produced. Low pressure laminating

techniques such as vacuum bagging and autoclaving used to fabricate

components directly to complex shapes can be expected to produce laminates

having an inherently larger variability.

Elastic Properties

Some recent data on the effect of cryogenic temperatures on Young's

modulus are presented on Figure 4 for the warp direction of fabric

laminates and for the edgewise direction of mat-reinforced laminates.

Where data are available, a distinction is made between the change

observed between 295 and 76 K and that observed between 76 and 4 K.

In contrast to strength data, the moduli data show a monotonical 1

y

increasing stiffness with decreasing temperature. Specific identification

of the materials is the same as for Figure 3.

The CERN data (9) provide the first cryogenic information on pitchbase fabric graphite/ epoxy laminates and illustrate the stiffness

superiority of this type laminate over conventional woven-fabric materials.

277
The high modulus is a direct reflection of the high modulus of the

graphite fiber. On the other hand, the CERN data indicate that the much

higher modulus of aramid fiber (124 GPa, 18 X 10 psi ) compared with

that of E glass (72 GPa, 10.5 X 10^ psi) is not reflected in the moduli

of woven-fabric laminates. The aramid-fabric laminates produced an

average Young's modulus very similar to that of the E-glass laminates

for equivalent fiber mass fraction.

Bars 1-4 of Figure 4 show that Young's modulus is directly related

to glass content. This again illustrates the significant effect of

glass content on laminate performance. The type of matrix has much less

effect because its modulus is typically an order of magnitude lower than

that of the glass reinforcement. These data indicate that the average

modulus of glass fabric-reinforced polymer laminates will increase about

20% upon cooling to 4 K.

The data presented on Figure 4 were obtained either from strain

gaged specimens or from extensometer data. NBS has also studied the

elastic properties of G-IOCR and G-llCR materials by dynamic resonance

methods (16-18). As illustrated on Figures 5 and 6, this method can

generate a continuous temperature-modulus curve for either the laminate

or the resin used in the laminate construction. Results, thus far obtained,

by the dynamic method agree reasonably well with those obtained by

conventional techniques. But the dynamic method has much improved

sensitivity. For example, the large difference in the temperature

dependence of Young's moduli of the G-IOCR and G-llCR resin systems

illustrated on Figure 6 is apparent in the differing temperature dependences noted for the G-IOCR and G-llCR warp direction laminate moduli

278
on Figure 5, but it is not apparent in data obtained from strain gaged

specimens. Figure 5 illustrates that the temperature dependence of

Young's modulus for the G-IOCR resin is substantially greater in the

direction normal to the reinforcement than in the warp direction.

Data presented on Figure 4 indicate that a polyimide-matrix laminate

(Bar 8) will have a Young's modulus very similar to that of an epoxy

matrix laminate of similar glass content (Bar 4).

The epoxy matrix laminate data on Figure 4 indicate a higher Young's

modulus than was reported for such laminates in the 1975 review (1).

This occurs because the earlier data reflected literature averages and

did not distinguish between moduli in the warp and fill directions. The

anisotropy of reinforcement results in anisotropy of Young's modulus as

illustrated for G-IOCR on Figure 7. These data suggest that average

moduli values reported for the E-glass fabric bisphenol A epoxy laminates

on Figure 4 would decrease about 40% if the moduli were evaluated in the

45® direction relative to the warp and fill of the fabric.

The Young's moduli of mat- reinforced laminates (bars 5 and 6 Figure

4) are substantially lower than those of fabric-reinforced laminates

when the latter are tested in the warp direction. However, the difference

decreases if the aforementioned orientation effect is considered. The

Young's modulus of a mat-reinforced product is essentially isotropic in

the plane normal to the pressing direction during laminate cure (edgewise di rection)

Elastic properties are i ntri nsi cal ly less variable than strength

properties.

279
Thermal Properties

Recent data on the temperature dependence of thermal expansion and

thermal conductivity of woven fabric and of random mat laminates are

summarized on Figures 8-10.

CERN (9) reports the first cryogenic thermal property data for

graphite and aramid-fabric reinforced epoxy laminates. Figure 8 shows

that the negative coefficients of thermal expansion of graphite and

aramid fibers combine with the positive coefficient of the epoxy resin

to yield laminates with a high degree of dimensional stability in the

warp direction (bars 10 and 11). The aramid laminate is more stable

than the graphite laminate of similar fiber content owing to the higher

negative coefficient of thermal expansion of the aramid fiber. The

stability of these products would be lower in the fill direction for an

unbalanced-weave fabric and would be still lower in the direction normal

to the reinforcement.

The thermal contraction of epoxy resins on cooling is about an

order of magnitude greater than for glass (19). Bar 12 on Figure 8

summarizes the contraction for a CY 205 + HY 905 + DY 040 + DY 061 resin

system evaluated by CERN. The data on Figure 8 indicate that the

contraction of a glass fabric-reinforced laminate is about 60% of the

contraction of the unreinforced resin in the normal direction and about

20% of the resin contraction in the warp direction. Contractions of the

aramid and graphite-reinforced laminates in the warp direction are only

about 5% of that of the unreinforced resin. This difference is reflected

in residual laminate stress. The comparatively small dimensional changes

observed on cooling from 76 K to 4 K indicate that most of the residual

stress created by cooling occurs in the 295 K to 76 K temperature range.

280
This stress is added to that created during cooling the laminates to

295 K from the elevated cure temperature. The low modulus of epoxy

(about 0.1 that of glass) results in most of the stress being resolved

onto the matrix.

The glass content would therefore be expected to play a significant

role in establishing contraction in the fiber direction. But such a

correlation is not evident in the data on Figure 8.

The thermal contraction o-f mat-reinforced laminates in the edgewise

direction is substantially higher than that for fabric-reinforced laminates

in the warp direction. This is due to the random fiber orientation and

lower glass content of the mat product. But contraction of the mat

products is isotopic in the edgewise direction while that of the fabric

product is anisotropic. Thermal contraction of fabric-reinforced laminates

in a 45° direction relative to warp and fill would probably be greater

than that of the mat-reinforced laminates in the normal direction.

Figures 9 and 10 summarize data reported on the temperature dependence

of thermal conductivity for a variety of wel 1 -characterized industrial

laminates (6, 9, 10, 11).

Figure 9 compares the conductivity of a series of E-glass reinforced

bisphenol A laminates and for two unreinforced bisphenol A epoxy resins.

The NBS data were obtained in the warp and normal directions on NEMA

G-10 and G-11 laminates produced to the CR grade specification (bars 1

and 2) and on the epoxy systems used for these grades (bar 3). Although

different resin systems are used for the two CR grades, their conductivities were sufficiently similar (within 5%) to justify their combination into a single bar on Figure 9 (20). CERN (9) characterized one

281
European commercial industrial laminate (Vetronite 432.86) in both the

warp and normal direction (bar 4) and a series of other commercial

industrial laminates in the normal direction. Five laminates had very

similar conductivites despite variations in glass content from 62-74

percent by weight (bar 5). One CERN laminate (Vetronite 541.02) displayed a significantly lower conductivity than the others (bar 6). CERN

also measured the thermal conductivity of a bisphenol A epoxy resin

system (CY 205 + HY 905 + DY 040 + DY 061) typical of that used in the

laminates (bar 7). Bar 8 reflects data obtained by a leading industrial

glass research company on a glass mat-reinforced epoxy product developed

for LNG use. Data on this product are available only to 76 K.

As a group, the NBS data reflect a much higher thermal conductivity

at 295 and 76 K than do the CERN data. Since the type and content of

reinforcement is about the same in all the laminates, it would appear

that the difference lies in the higher conductivity of the resin used in

the NBS laminates (bar 3) compared that used by CERN (bar 7). A comparison of the data on Figure 9 with previously published literature

averages for thermal conductivity of glass fabric-reinforced laminates

(1) shows that such averages correlate well with the CERN data. Work to

clarify this apparent discrepancy is presently underway at NBS.

Figure 10 compares the CERN data (9) on the temperature dependence

of thermal conductivity for E-glass fabric-reinforced silicone (Vetronite

432.91) and polyimide laminates (Silirite 64 400, Imelite 6 101) and for

aramid (Kevlar 49) and graphite-fabric (Thornel PVCB 45) reinforced

bisphenol A epoxy laminates. All data are for the direction normal to

the reinforcement.

282
The comparatively high thermal conductivity of the graphite-fabricreinforced laminate (bar 5) is due to the relatively high conductivity

of graphite compared to glass. The E-glass reinforced polyimide laminates

(bars 2 and 3) have conductivities very similar to those reported by

CERN (9) for epoxy matrix laminates (cf Bar 5, Fig. 9). This suggests

that the thermal conductivity and the temperature dependence of conductivity are very similar for polyimide and epoxy resins.

The normal direction thermal conductivity for the glass/silicone

and aramid/epoxy laminates are substantially lower than for glass/epoxy

laminates of equal fiber mass fraction. Normal direction conductivity

for cotton phenolic laminates ( phenol formal dehyde resin) is similar to

that for glass/epoxy laminates.

The author is unaware of any substantial contribution to the

characterization of specific heat properties of fabric or mat-rei nforced

laminates within the cryogenic range since the earlier reviews (1,2).

Heberlein (21) has published data indicating that the specific heat of

E-glass laminates might be substantially lower than that for laminates

fabricated with S glass. But these data did not extend below 230 K.

Electrical Properties

The effect of cryogenic temperatures on the volume resistivity and

electrical strength of G-IOCR and G-llCR glass fabric-reinforced epoxy

laminates is summarized on Table I (6). These are the first data obtained

on these parameters at 4 K. Although the data are not extensive, they

indicate that the electrical strength will be independent of both temperature and material. The effect of temperature on the volume resistivity

is similar for the two materials, increasing 1.5-2 orders of magnitude

on cooling from 295 K to 4 K.

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UNIAXIAL FIBER-REINFORCED COMPOSITES

OVERVIEW

Composites reinforced with uniaxial glass, graphite, boron or

aramid fibers continue to have limited and specialized applications in

cryogenic technology. The aerospace industry is using graphitereinforced laminates for dimensionally stable mirrors and reflectors

(4). The relatively small structures already flown have functioned well

in the temperature range down to 100 K. Primary concern is predictability

of dimensional stability in the very large reflectors presently being

designed for use in outer space.

The desire to reduce the weight of superconducting machinery and of

superconducting magnetohydrodynamic power generators has provided the

primary impetus for continued study of the performance of advanced

composites at cryogenic temperatures. Most of the related research has

been carried out by NBS with the assistance of U.S. industry (22-25)

and by CERN (9).

The 1975 review (2) indicated that uniaxial glass, graphite, boron

and aramid fibers were all viable reinforcements for cryogenic laminates.

It was shown that a wide variety of polymer matrices could be used,

although epoxy systems emerged as the favorite choice. A primary problem

was the large variability in the mechanical property data reported in

the literature. Unfortunately, it was not possible to distinguish

inherent material variability from experimental variability. The experimental problems are now under better control, although much remains to

be done in that area. Improved test methods have established that the

earlier indications of strength degradation in graphite-reinforced epoxy

234
laminates on cooling were incorrect. Further work has also shown that

the use of highly flexibil ized resin systems is counter productive in

the advanced composites as well as in the gl ass-rei nforced products.

Conventional, fully reacted resin systems are the best current choice

for high-modulus fiber laminates intended for cryogenic use.

Recent characterization studies on advanced composites have produced

mechanical property data of the type required for prediction of the

properties of complex angleply laminates using macromechanical laminate

theory (26,27). These data, obtained from uniaxial laminates, consist

of longitudinal and transverse uniaxial strength, modulus and Poisson's

ratio plus in-plane shear strength and modulus. Laminates reinforced

with S glass; high, medium and low modulus graphite; boron and aramid

fibers have now been characterized in this manner. Results of these

studies have been published in NBS reports (24) and in the open literature

(22,23,25).

DATA ANALYSIS

The most recent data on the temperature dependence of the significant

mechanical and thermal properties of uniaxial fiber reinforced composites are summarized on Figures 11-17. Bar graphs are again used to

illustrate trends of strength changes from room temperature to the

cryogenic temperature range for data obtained at fixed cryogenic boiling

points. The average of data reported at 76 K and 4 K are plotted

except where the data show a statistically significant change in a

property between these temperatures . Data obtained at closely-spaced

temperature intervals are presented as continuous curves.

235
Strength Properties

The uniaxial longitudinal and transverse tensile strengths of

glass-, boron-, graphite- and aramid-reinforced epoxy laminates reported

by NBS (22-25) and by CERN (8) are summarized on Figure 11.

Data for epoxy matrix laminates uniaxially reinforced with three

types of glass appear on Fig. 11. Typical compositions of each glass

are given on Table II.

Type S glass was developed during the early 1960's as a highstrength glass for specialized aerospace purposes. The 901 designation

refers to a special solvent-sizing system used with this high-cost,

military-specification product. The use of S glass is confined to

filament-wound or uniaxial prepreg tape products. Most of the cryogenic

data on uniaxial glass-reinforced epoxy laminates have been generated

with Type S glass. But potential users should also consider the recently

introduced commercial grade of S glass produced under the trade designation

S-2. For reportedly comparable room temperature properties, type S-2

is a much lower cost product (28). There are presently no cryogenic

data for laminates fabricated from S-2 glass. Type S-2 glass is available

in woven fabrics as well as uniaxial fiber.

Type E is the standard electrical insulating grade of glass, most

often used in fabric-reinforced industrial laminates but also available

in uniaxial form. It is substantially less expensive than types S or S2. Type R is a glass of European manufacture used in laminates evaluated

by CERN (9).

These data show that the uniaxial tensile strength of the S-901

glass/Resin 2 laminate averages approximately 30% higher than that of

286
the E-gl ass/bi sphenol A laminate at room temperature and about 50%

higher at cryogenic temperature. This difference is consistent with

published data on the comparative strengths of the two glass types (28).

The R-gl ass/bi sphenol A laminate developed uniaxial tensile strengths

intermediate between that of the S-901 and E-glass laminates.

Uniaxial longitudinal tensile properties are expected to be fiber

dominated. But there is evidence in the literature that the average

strength reported for the S-901 laminate on Fig. 11 has been substantially

lowered by the use of the highly-flexibil ized Resin 2 epoxy formulation

Handbook data (29) suggest that the room temperature tensile strength of

this laminate is about 30% lower than expected from an S-901 glassreinforced laminate of equal glass content in a conventional, fully

reacted epoxy matrix. Cooling has increased the strength of the S901/Resin 2 laminate, but only to a value expected for a conventional

matrix laminate at room temperature. A similar discrepancy between the

strengths of S-901 -glass reinforced laminates involving the two types of

matrices was documented in a review article published by the author in

1975 (3), but the implication was not recognized at that time.

The useful strength of glass-reinforced laminates at cryogenic temperatures will be much lower than the ultimate strength whenver damage

accumulation in the matrix is unacceptable. Uniaxial glass-reinforced

laminates have failure strains approaching 3% at 4 K; however, the

embrittled matrix will typically fail at less than 1% strai n--about

20-30% of the ultimate fracture stress. The stress level corresponding

to the onset of matrix damage may therefore be lower at cryogenic temperatures than at room temperature. This factor makes Type S glass less

287
functionally attractive compared to Type E glass for applications where

the increased damping and the slightly decreased modulus accompanying matrix damage cannot be tolerated. Similar functional strength

limitations do not exist for laminates reinforced with high-modulus

fibers such as graphite or boron for which the ultimate strain to failure

is less than 1% at all temperatures.

Figure 11 indicates that the graphite-reinforced laminates are much

better cryogenic materials than was indicated by data available in

1975 (2). The latter data, compiled from literature sources, suggested

that cooling would lower the tensile properties of graphite-reinforced

epoxies as a group. Correlation between laminate and fiber properties

was not observed and a large scatter in temperature dependency of

strength properties was reported. Figure 12 shows that this behavior is

not confirmed by the more recent studies at NBS and CERN. It now appears

that the effect of cryogenic temperatures on the average uniaxial

longitudinal tensile properties is minimal for these materials.

The properties of the matrix and of the matrix-fiber bond strongly

influence the uniaxial transverse tensile strength. Thus, the very high

transverse strength illustrated for boron-reinforced aluminum on

Figure 11 reflects the strength of the 6061 aluminum alloy matrix. The

low transverse strength observed for the AS/Resin 2 graphite laminate

at 295 K reflects the weakness of this highly-flexi bil ized resin at

higher temperatures, while the significant decrease in strength on cooling

is attributable to a high state of internal stress arising from the high

thermal contraction of this resin system combined with the negative

238
coefficient of thermal expansion of the graphite fiber. Work at NBS

(24) has shown that the average transverse strength of an unfl exi bil ized

,

conventional, high-strength graphite laminate averages 48 MPa(7 X 10 psi)

at 4 K; over 300% higher than that developed by the flexibil ized-matrix

laminate.

The tensile data for unreinforced bisphenol A epoxy illustrated on

Figure 11 is reported by CERN for a CY 205 + HY 905 + DY 040 + DY 061

system. The data should be typical of that developed by conventional,

fully reacted epoxy resin systems. Lower transverse tensile strengths

of the laminates compared to the resin strength indicates low interfacial bond strengths in the laminates.

Recent data on the temperature dependence of the uniaxial longitudinal

tensile strength of aramid-fi ber-rei nforced epoxy appearing on Figure 11

is in agreement with prior literature values (2). These data and the

newly acquired data on the uniaxial transverse tensile strength appear

to be temperature independent.

The aramid laminate data appearing on Figure 11 reflect average

properties of a conventional laminate system. The large negative coefficient

of thermal expansion of aramid fiber has led to the belief that the

cryogenic performance would be improved if a highly flexibilized matrix

such as the Resin 2 formulation was used to accomodate the differential

contraction. But NBS studies of the performance of an aramid-fiber

laminate using the Resin 2 system has shown that its use is counterproductive.

The transverse tensile strength of this combination was found to be so

low (0.3 MPa, 42 psi at 245 K and 0.1 MPa, 17 psi at 76 K) as to render

the laminate almost useless for structural applications at any temperature.

2J9
Again, it is believed to be the weakness of the resin at 295 K and the

development of a high state of internal stress at cryogenic temperatures

that destroys the performance of the laminate.

Reliable compressive strength properties are difficult to obtain

in composite materials. A primary difficulty is the strong dependence of

measured values on specimen alignment and on proper specimen support.

The large surface-to-vol ume ratio of typical specimens accentuates the

significance of surface flaws. Particularly for the uniaxial longitudinal mode, values tend to have a distribution function which is

highly skewed toward low values. The magnitude of the scatter increases

as the compression strength increases. Average values of compression

data obtained from laboratory tests therefore tend to be systematical ly

biased toward lower values than the true intrinisic material strength.

The maximum compression strength values obtained in a test series may

more accurately reflect the true compression strengths of high-strength

composites than does the mean of the distribution. Uniaxial transverse

compressive strength data show comparatively little scatter and a relatively

normal distribution. Here, the distribution mean is perhaps the most

meaningul indication of compression strength.

Figure 13 summarizes the most recent data obtained by NBS (23-24)

on the effect of cryogenic temperatures on compressive strengths. A

comparison is made between the uniaxial longitudinal and transverse

compressive strengths of epoxy laminates reinforced with glass, boron,

graphite and aramid fibers and aluminum reinforced with boron. All data

were obtained by an unsupported beam test method.

290
Results suggest that the intrinsic uniaxial longitudinal compressive

strength of typical (0.14 mm, 5.6 mil diameter) boron-reinforced laminates

is on the order of 3450 MPa (500 X 10 psi ) at cryogenic temperatures.

Differences in strength noted for the epoxy and aluminum-matrix products

are attributable to the noted differences in fiber content. These data

are at variance with prior data (2) which had indicated a negligible

temperature dependence of compressive strength for boron-reinforced

aluminum. The recent data also indicate a much higher compressive

strength than previously reported in the literature for unsupportedbeam tests. The data on Figure 13 approach calculated compressive

strengths based on sandwich-beam specimens (30). Comparative room

temperature data were not available for these boron-reinforced materials.

The boron-reinforced materials are clearly the strongest laminates

in the transverse as well as in the longitudinal compression. The

transverse compressive strength of the 5.6-mil boron-aluminum laminate

is substantially better than that reported in 1975 (2). Cooling from

76 to 4 K improves the transverse strength of the aluminum-matrix

material while having little effect on the epoxy-matrix laminates. This

is due to the substantial (20-40%) increase in strength of the 6061

aluminum matrix between these two temperatures (31).

S-901-glass laminates develop compressive strengths about half that

of boron laminates, followed in strength by the graphites. Aramid laminates have the lowest compressive strength. But the use of the highly

flexibilized Resin 2 epoxy matrix appears to have reduced the compressive properties of the S-901 and type AS graphite laminates compared

with similar laminates using conventional epoxy resin matrices. The
room temperatures values reported for the S-901/Resin 2 laminate on

Figure 13 are about 20% lower in the longitudinal direction and about

50% lower in the transverse direction than that reported for a similar

laminate using a conventional resin system (3M XP251S) in MIL-HDBK1 7A (14). As was the case with the tensile properties of this material,

cooling substantially improved the compressive strength of the S-901/Resin 2

laminate at cryogenic temperatures. But improvement in the transverse

direction is only sufficient to return the laminate to a transverse

compression strength expected from a conventional laminate at room

temperature. The deleterious effect of using the highly flexibilized

epoxy resin was also seen in the low strengths developed in the laminate

reinforced with type AS graphite fiber. Despite a high inherent fiber

strength, the longitudinal compressive strength of this laminate has

been reduced to no higher values than that obtained with the much lower

strength GY-70 fiber laminate. Furthermore, a significant reduction in

longitudinal compression strength is noted for the AS/Resin 2 laminate

on cooling from 76 to 4 K, a phenomenon not observed in the other laminates.

A deterioration in compressive strength is also evident in the transverse

direction for this laminate.

The comparatively low longitudinal compressive strength of the

aramid-fiber laminate reflects low compressive fiber properties. The

strength is low both in comparison to the other type of laminates and as

a percentage of its uniaxial tensile strength. The difference is not as

noticeable in the transverse compression test mode where strength is

more influenced by matrix properties.

292
Elastic Properties

Figure 14 summarizes the data obtained by NBS (22-25) and by CERN (9)

on the temperature dependence of Young's modulus in tension for uniaxial

longitudinal and transverse directions of glass-, boron-, graphiteand aramid-rei nforced epoxy laminates.

Only the aramid-reinforced epoxy laminate displayed a significant

temperature dependence in the longitudinal direction. Present values

the aramid laminate are significantly lower than values reported for

this type laminate in the earlier literature review (2). Furthermore,

the substantial modulus increase reported to occur in aramid-reinforced

epoxy laminates below 76 K (32) was not observed by NBS.

The longitudinal tensile moduli of glass-reinforced laminates are

dominated by the moduli of the glass which is about an order of magnitude

higher than that of the conventional bisphenol A epoxy system illustrated

on Figure 14. The longitudinal modulus of a glass-reinforced laminate

is therefore strongly dependent on the volume fraction of glass in the

laminate. It is also influenced by the type of glass reinforcement. The

modulus of type E glass is 10-20% lower than that of the S type (28).

The high transverse modulus reported for the S-901 -glass laminate

suggests that the type of glass, and perhaps the finish on the glass,

has a strong influence on modulus in this direction.

The S-901 -glass laminate displays a large percentage increase in

transverse modulus on cooling. This reflects hardening of the highly

flexibilized Resin 2 epoxy matrix in this material. However, a comparison of the present S-901 laminate data with handbook values (29)

293
shows that the Resin 2 laminate must be cooled to cryogenic temperatures

before developing a modulus equal to that produced at room temperature

by a conventional S-glass laminate using a nonfl exi bil ized , fully

reacted epoxy resin system.

The boron-reinforced laminates provide an excellent combination of

longitudinal and transverse tensile moduli. Doubling of the transverse

modulus of the boron/epoxy material is the only significant temperature

effect.

The Young's modulus in tension of the AS/Resin 2 laminate is 20-25%

lower in both the longitudinal and transverse direction than expected

for a laminate reinforced with this type of fiber in a conventional,

fully reacted epoxy matrix (30).

The temperature dependence of inplane shear moduli summarized on

Figure 15 represents the first cryogenic data obtained on this parameter.

These data, obtained by the ±45° tensile coupon method (33), indicate a

substantial increase of average shear modulus for all types of laminates

studied. In-plane shear modulus is strongly influenced by the modulus of

matrix. Values reported for the S-901 -glass and type AS-graphite laminates

on Figure 15 are therefore lower than would be expected in conventional

laminates due to the low room temperature shear strength of the highly

fl exi bil ized Resin 2 epoxy formulation used in their farbi cation.

Thermal Properties

Recent data on the effect of cryogenic temperatures on dimensional

stability and thermal conductivity of uniaxial laminates are summarized

on Figures 16 and 17.

294
Figure 16 summarizes data on the effect of temperature on dimensional

stability of well characterized epoxy-matrix laminates reinforced with

glass, boron, and graphite and of boron-reinforced aluminum. Data are

also provided on the stability of a typical, fully reacted bisphenol A

epoxy resin system and on the highly flexibilized Resin 2 epoxy resin

system used with several of the laminates.

The thermal contraction of the laminates is much less than that of

the matrix materials. Particularly in the fiber (longitudinal) direction,

the reinforcement properties dominate the contraction because the reinforcement modulus is much higher than that of the matrix. The longitudinal

modulus is therefore dependent on the fiber volume fraction as well as

on the type of fiber, with the matrix playing a minor role.

Thermal stability in the transverse direction of uniaxial laminates

is influenced by both matrix and reinforcement properties. A further

analysis of the data of Figure 16 shows that the transverse contraction

of the glass- and graphite-reinforced bisphenol A laminates (fiber types

R, E, HM-S) was about 52% of that expected for the unreinforced bisphenol A

resin, whereas the glass- and graphite-reinforced Resin 2 laminates (S901, AS) contracted about 35 % of that expected from unreinforced Resin 2.

The lesser contribution of the Resin 2 matrix to the laminate contraction

is due to the low modulus of this highly flexibilized resin at temperatures

above the cryogenic range. Comparable values for the boron/epoxy and

boron/aluminum laminates are 32 % and 63%, respectively. Here, the

comparatively low transverse contraction of the epoxy laminate is due to

the very high modulus of the boron compared with the epoxy. The aluminum

contributes more to determining the contraction of the laminate because

its modulus is much higher than that of epoxy.

295
The temperature dependence of thermal conductivity for uniaxial

laminates in the longitudinal and transverse direction are summarized on

Figure 17. Data for an unreinforced conventional bisphenol A epoxy is

included for comparison. All of the epoxy-matrix laminates used a

conventional epoxy system. The boron/aluminum laminate had a 6061

aluminum alloy matrix.

Uniaxial laminates are most frequently used as thermal standoffs in

cryogenic technology because of the low conductivity of the epoxy-matrix

products combined with excellent strength properties. Glass-reinforced

epoxy composites are favored because, as illustrated on Figure 17, they

provide lower conductivity in the fiber direction than boron, graphite,

or aramid-rei nforced composites. Boron and aramid-reinforced laminates

appear to have about the same conductivity in the longitudinal direction--

approximately 20% higher than that of the glass laminate. The graphitereinforced epoxy laminate has a longitudial conductivity about 45%

higher than that for glass reinforcement. The very high conductivity of

the aluminum-matrix laminate limits its use to applications having

minimum thermal gradients or to specialized applications where a high

thermal conductivity is desired.

Transverse conductivities are all lower than in the longitudinal

direction regardless of the relative conductivities of the fiber and

matrix. This reflects the thermal resistance of the fiber-matrix interface in the transverse direction.

The aramid-reinforced epoxy laminate displays the largest change in

thermal conductivity on cooling. The transverse conductivity data

296
plotted for this material on Figure 17 were measured in the direction

normal to the rei nforcement of a fabric-reinforced laminate (9). It is

therefore only an approximation to the transverse conductivity of a

uniaxial laminate. But the approximation should be quite good for

similar fiber contents in the two types of laminates.

FUNDAMENTAL RESEARCH

When composites first became of interest for cryogenic applications

during the space programs of the 1960's, it was commonly believed that

the primary problem would be the embrittling effect of low temperatures

on the polymer systems used for laminate matrices. Research (34) led to

the development of an epoxy system known as Resin 2*. The unreinforced

resin provided improved shock resistance and low temperature strain

capability through the addition of fl exi bil izing agents. This system

was intended for filament winding of cryogenic pressure vessels. The

maximum filament stress attainable at cryogenic temperatures was

reportedly improved with the Resin 2 formulation compared with that

attainable using conventional epoxy systems (34). But, as is clear from

the data summarized in preceeding sections of this review, extending the

use of the Resin 2 formulation to glass-, graphite- and aramid-fiber

reinforced laminates fabricated by other means is counter productive.

The mechanical performance of laminates using Resin 2 epoxy is distinctly

inferior to laminates fabricated with conventional, fully reacted epoxy

resin systems at both cryogenic or room temperature. It now appears

that room temperature performance is degraded by the extreme weakness of

the resin and cryogenic performance is degraded by high levels of

residual stress created by the large thermal contraction of the Resin 2

*Epon 828/DSA/Empol 1040/BDMA in proportions 100/115.9/20/1 pbw.

297
systems compared with that of conventional resin systems. Such stresses

substantially limit the effective strain capability of the laminates at

4 K.

Conventional, fully reacted epoxy resin systems are therefore the

best current choice for the matrices of laminates to be used at cryogenic temperatures . But the need to develop improved matrix performance

for cryogenic laminates remains. The range of useful stress obtainable

at cryogenic temperatures is limited in conventional glass-reinforced laminates

by the onset of damage in the embrittled matrix at stresses on the order

of 20-30% of the ultimate laminate stress. This is lower than the

useful stress at room temperature, despite the much higher ultimate

fracture stress at low temperature. Extension of this useful range by

improving the effective cryogenic strain capability of the laminate

matrix is very desirable. An improved matrix resin would also reduce

the microcracking observed in some angleply graphite-reinforced laminates

subjected to thermal shock.

But the primary benefit expected from such a resin system is reduced

variability in cryogenic strength properties. A major contributor to

the variability displayed by conventional laminates is the decrease in

toughness and increase in notch sensitivity of fully reacted epoxy resin

systems when cooled to cryogenic temperatures. This causes the laminates

to become very sensitive to the presence of flaws. In fracture mechanics

terms, the critical crack size becomes extremely small. An illustration

of this effect is provided by Figure 18 , which summarizes the range of

uniaxial longitudinal compressive values obtained during NBS testing of

a series of graphite-, glass- and boron-reinforced laminates at 295 K,

76 K and 4 K. The data have been normalized at 295 K to illustrate

differences in the magnitude of the data spread for each material upon

298
cooling. A positive correlation is noted between the amount of data

scatter and the average increase in compression strength on cooling.

Comparison with Figure 13 shows a positive correlation also exists

between variability and inherent compressive strength of the materials.

This illustrates that a flaw in the embrittled matrix has increasing

significance as compression strength increases. Variability attainable

to the increasing notch sensitivity of the resin matrix is a particular

problem in compression, shear and off-axis tension properties.

Possibilities for improving epoxy systems for use in cryogenic

composite laminates have been recently addressed by Kreibich, Lohse, and

Schmid (35) and by Hartwig (36). The possibility of achieving cryogenic

flexibility by modifying epoxy resins to facilitate rotational modes of

long-chain molecules (crankshaft rotation) is being considered. Such a

system would theoretically retain good room temperature strength, have a

relatively low thermal contraction, and provide improved cryogenic

toughness. Frank (37) has also contributed to the fundamental understanding

of polymer embrittlement by considering the role that far infrared

spectroscopy can play in assessing the degree of crystallinity existing

at cryogenic temperatures. Schmid (38) has considered the connection

between secondary valence forces, primary valence energies and low

temperature flexibility of polymers.

Kneifel (39) has described some of the initial work undertaken at

the Institut fur Experimentel 1 e Kernphysik, Karlsruhe, Germany, on the

fracture properties of epoxy resins at low temperatures . This work has

shown that the fracture toughness, Kj^, of rigid and semi flexible resins

is much higher at 76 K than at 5 K, indicating the presence of relaxation

299
processes at 76 K. A substantial difference was noted in the 5 K

fracture toughness among rigid, semi flexible and flexible epoxy resin

systems at room temperature. It is perhaps significant that the flexibilized resins did not show the highest 5 K toughness.

The conference on Nonmetal lie Materials and Composites at Low

Temperatures held in Munich, West Germany July 10-11 , 1979 (4, 16,

35-39) was a valuable forum for technology transfer in these fundamental

areas. Additional conferences addressing this subject are planned by

the International Cryogenic Materials Conference (ICMC).

300
TEST METHOD DEVELOPMENT

Variability of reported data on the mechanical performance of

composite materials at cryogenic temperature has been identified in this

paper as a significant concern. The variability arising from the experimental

procedure must be minimized to allow inherent material variability to be

assessed.

Cryogenic temperatures increase the difficulty of obtaining valid

test results. The cost of cryogenic testing tends to put practical

limits on the number of specimens used to evaluate a parameter. Furthermore,

the high surface- to-vol ume ratio of typical laboratory specimens enhances

the probability of premature failure due to the embrittling effect of

cryogenic temperatures on the resin matrix.

Techniques for obtaining valid fractures within the gage length

sufficiently remote from grip effects during uniaxial longitudinal

laminate testing have been reasonably well solved at NBS (22-23) and at

CERN (40). The technique is to design a grip system permitting some

Poisson's contraction to take place during the test within the portion

of the grip nearest to the gage section, thereby minimizing stress

concentrations in that area. Compression test fixtures which work well

at 4 K have been developed at NBS (23) and refined at CERN (40). Problems

of minimizing liquid helium consumption, of achieving proper specimen

alignment and of achieving proper specimen support are fairly well in

hand, although refinements will continue to be made.

Some applications of composite materials in superconducting technology place unique demands on laminate interlaminar shear strength.

Grooved sections of glass-fabric reinforced epoxy laminates are frequently

used to support and insulate superconducting windings. The laminates

301
must withstand electromagnetic forces resolved in shear across the

portion of the laminate separating the windings. Attempts by NBS and

MIT to assess the effect of cryogenic temperatures on interlaminar shear

strength have shown results with ASTM test method D 2733-70 ("guillotine"

method) are highly variable. Results depend, to a large extent, on how

the method is applied in practice. This method, which assesses the

stress to separate adjacent plies located between opposing offset slots

machined in to the specimen surface, was found to be highly dependent on

the degree of specimen side support and on the ratio of slot separation

distance (L) to specimen thickness (T). Becker and Erez (15) found that

interlaminar strength values increased continuously as the L/T ratio of

the specimen decreased, approaching that measured by the torsion method

at L/T = 0. Kasen et al.(6) found systematic increases in interlaminar

shear strength by systematical ly increasing side support pressure for a

specimen having an L/T of 2. A limiting strength level about equal to

that obtained in an unsupported specimen at an L/T value of about 0.83

was defined.

Additional refinement of test methods for assessing interlaminar

shear is clearly required. But there is some consolation in observing

that inadequacies of existing test procedures are probably yielding data

which are conservative for design purposes.

The thermal conductivity data presented on Figure 9 indicate the

need for refining methods of measuring the conductivity of polymermatrix laminates in the higher temperature ranges. It is presently

necessary to introduce large correction factors into the measured values

302
to account for radiation losses from semitransparent laminates such as

glass-reinforced epoxy. The effect of specimen thickness on validity of

the data must also be more precisely evaluated. These factors are

currently being assessed at NBS.

A neglected test-related factor is the specimen size effect. This

may be more of a problem in evaluating mechanical property values for

composite materials at cryogenic temperatures than at room temperatures

due to the increased notch sensitivity of polymer matrices at low temperatures.

Dynamic resonance techniques for assessing the temperature effect

on elastic constants of composites have been under development at NBS

(16-18). Results of this approach presented on Figures 5-7 illustrate

the flexibility it offers in assessing elastic constants at arbitrarily

selected directions through a laminate. A comparison of Young's moduli

values obtained by dynamic methods with those obtained by conventional

strain-gaged specimen methods shows good correlation for a number of

laminate types. But dynamic methods cannot discriminate between differences

in tensile and compressive modul i--resul ts are always averages. There

is also some concern that the very low levels of strain inherent in conventional

dynamic methods permit sensing of only first-order strain terms in some

fabric-reinforced laminates. In this case, moduli values will differ

significantly from those derived by mechanical means which incorporate

additional strain terms significant in defining practical engineering

laminate moduli. Phenomenologically, the usual dynamic modulus methods

examine only the very initial portions of a stress-strain curve where

the slope may or may not be the same as for larger strains. This problem

303
can be overcome by applying dynamic methods to specimens having desired

levels of prestrain. Studies in this direction are under way at NBS.

Dynamic methods produce data on damping as well as elastic properties,

providing a convenient method of following damage accumulation in

laminates subjected to cyclic mechanical stress or to thermal fatigue (41).

304
SUMMARY AND RECOMMENDATIONS

Recent studies clarifying the effect of cryogenic temperatures on

the mechanical and physical performance of fiber-reinforced composite

materials presented in this review confirm that conventional, off-theshelf laminates developed for room temperature service are viable engineering

materials to temperatures as low as 4 K. Averages of static mechanical

properties and elastic properties improve on cooling. But it is becoming

clear that embrittling of the polymer matrix and the consequent decrease

in strain capability is limiting the extent to which effective use can

be made of such improvements. Many glass-reinforced laminates are in

practice limited to 20-30% of their ultimate tensile strength at cryogenic

temperatures because of damage accumulation in the matrix at higher

stress levels. This occurs because cooling decreases the strain capability

of the polymer matrix while increasing that of glass. Matrix embrittlement

does not limit the useful stress range of laminates reinforced with high

modulus fibers such as graphite or boron because of the inherently low

elongation to fracture of the fibers. But for graphite laminates it

appears that the negative coefficient of thermal expansion of the fibers

combines with the positive coefficient of the polymer matrix to cause microcracking in the embrittled polymer. Embrittlement of the matrix reduces

toughness, creates notch sensitivity and increases the variability in

mechanical properties at cryogenic temperatures, despite substantially

increased average property values.

It therefore seems appropriate to define the primary technical

problem as embrittlement and decreased toughness of existing polymer

systems used for composite matrices. The approach being taken by

305
Ciba-Geigy (35) and the Institut fur Technische Physik, Karlsruhe, West

Germany (36) are therefore most appropriate. This work should be encouraged

along with studies designed to more clearly associate the failure modes

and life-cycle performance of laminates with their constitutive properties and the geometry of laminate construction.

ACKNOWLEDGMENTS

This review was sponsored by the USAF Aero Propulsion Laboratory,

Wright-Patterson Air Force Base, Dayton, Ohio under contract FY1455-79-

00603. Additional support was provided by the Office of Magnetic Fusion

Energy, Department of Energy under Interagency Agreement EA 77-A-Ol6010.

The author is particularly indebted to Dr. Knud Dahl rup-Peterson of

CERN for making his data available, for consultation and for review of

the final manuscript. The contribution of data and consultation of Dr.

John Olinger of the Owen/Corning Research Laboratory is also gratefully

acknowl edged.

306
PROPRIETARY MATERIALS

The following proprietary materials referred to in this report are those used by the cited author. Generic names have been substituted

wherever it was possible to do so without sacrificing clarity. The use of trade names in no way implies approval, endorsement or recommendation

of specific commercial products by NBS

Fibers and Reinforcements:

S, S-901 , S-2 Owens/Corning Glass Company

Kevlar 49 E. I. Du Pont de Nemours & Company

AS, HM-S Courtalds Ltd.

GY-70 Celanese Corporation

R Saint-Gobain Industries

Resin System Components

CY 205, HY 905, DY 040

DY 061

Ciba-Geigy AG

Empol 1040 Emery Industries, Inc.

Epon 828 Shell Chemical Company

DSA dodecenyl succinic anhydride

BDMA

Lami nates

benayl dimethyl ami ne

Mi carta Westinghouse Corporation

Permaglas T630

Vetronite 54.02, 432.81,

432.86

Durostone EPF 2, EPF 4

Si li rite 64 400

European commercial laminates

307
REFERENCES

1. Kasen, M. B. "Mechanical and thermal properties of filamentaryreinforced structural composites at cryogenic temperatures, 1:

glass-reinforced composites," Cryogenics 15 (1975) 327.

2. Kasen, M. B. "Mechanical and thermal properties of filamentaryreinforced structural composites at cryogenic temperatures, 2:

advanced composites," Cryogenics 15 (1975) 701.

3. Kasen, M. B. "Properties of fil amentary-reinforced composites at

cryogenic temperatures" ASTM STP 580, American Society for Testing

and Materials (1 975) 586.

4. Kasen, M. B. "Cryogenic applications of composite technology in the

U.S.A. Nonmetallic Materials and Composites at Low Temperatures

(1 979) Plenum Press, NY 317.

5. Kasen, M. B. "Composite laminate applications in magnetic fusion

energy magnet systems" Proceedings 1978 International Conference

on Composite Materials (1978) 1493.

6. Kasen, M. B., MacDonald, G. R. , Beekman, D. H. and Schramm, R. E. "Mechanical, electrical and thermal characterization of G-IOCR and

G-llCR glass cloth/epoxy laminates between room temperature and

4 K" Advances in Cryogenic Engineering , V 26 Plenum Press, NY

(1 980)

7. Kernohan, R. H., Coltman, R. R. and Long, C. J. "Radiation

effects on organic insulators for superconducting magnets,"

Oak Ridge National Laboratory Report ORNL/TM-6708 (1979).

8. Coltman, R. R. , Klabunde, C. E., Kernohan, R. H. and Long, C. J. "Radiation effects on organic insulators for superconducting magnets," Oak Ridge National Laboratory Report ORNL/TM-7077 (1979).

9. Dahl rup-Peterson , K. and Perrot, A. "Properties of organic com- posite materials at cryogenic temperatures: CERN ISR-BOM/79-39."

(November, 1979).

10. Olinger, J. Owens/Corning Fiberglass Corporation Technical Center,

Granville, Ohio (personal communication).

1 1 . LNG Materials and Fluids-A Users Manual of Property Data in Graphic

Fo rmat. Supplement 2 to First Ed. (1 980). Cryogenic Data Center,

National Bureau of Stnadards, Boulder, CO 80302.

12. Stone, E. L. , El-Marazki, L. 0. and Young, W. C. "Compressive

fatigue tests on a unidirectional glass/polyester composite at

cryogenic temperatures" Nonmetallic Materials and Composites at

Low Temperatures Plenum Press, NY (1979) 283.

308
13. Stone, E. L. and Young, W. C. "Coefficient of friction measurements

of fiberglass/epoxy at cryogenic temperatures" Advances in Cryogenic

Engineering V 26 Plenum Press, NY (1980).

1 4 . Plastics for Aerospace Vehicles, Part 1. Reinforced Plastics.

MIL-HDBK-1 7A, Department of Defense, Washington, D.C. (Jan. 1971).

15. Becker, H. and Erez, E. "A study of interlaminar shear strength at

cryogenic temperatures" Advances in Cryogenic Engineering V 26

Press NY (1980).

16. Ledbetter, H. M. "Dynamic elastic modulus and internal friction in

fibrous composites" Nonmetallic Materials and Composites at Low

Temperatures Plenum Press NY (1979) 267.

17. Ledbetter, H. M. "Dynamic elastic modulus and internal friction in

G-IOCR and G-llCR fiberglass-cloth-epoxy laminates" NBS Technical

Report NBSIR 79-1609 (1979) 437.

18. Ledbetter, H. M. in "Materials Studies for Magnetic Fusion Energy

Applications at Low Temperatures: II," Eds, F. R. Fickett and

R. P. Reed, NBSIR 79-1609 (1979) 437.

19. Corruccini, R. J. and Gniewek, J. J. "Thermal expansion of technical

solids at low temperatures; A compilation from the literature."

NBS monograph No. 29 (1961).

20. Must, J. NBS, Boulder, Colorado (personal communication).

21. Heberlein, D. C. "Heat capacities of E-glass, cross-ply; S-glass,

cross-ply; and S-glass, unidirectional epoxy composites" U.S. Army

Mobility Equipment Research and Development Command Report 2199,

Fort Bel voir, Virginia (1 976).

22. Schramm, R. E. and Kasen, M. B. "Static tensile properties of

boron-aluminum and boron-epoxy composites at cryogenic temperatures"

Advances in Cryogenic Engineering V 22 Plenum Press, NY (1977) 205.

23. Schramm, R. E. and Kasen, M. B. "Cryogenic mechanical properties of

boron-, graphite- and glass-reinforced composites" Materials Science

and Engineering 30 (1977) 197.

24. Schramm, R. E. and Kasen, M. B. "Semi-annual report on materials

research in support of superconducting machinery: advanced composites"

NBS-ARPA Reports NBSIR 74-359 (1974) 101, NBSIR 74-393 (1974) 123,

NBSIR 75-810 (1975) 123, NBSIR 75-828 (1975) 1, NBSIR 76-839

(1 976) 1 , NBSIR 76-848 (1976) 1 .

25. Kasen, M. B., Schramm, R. E., and Tobler, R. L. "Advanced composites

for cryogenic applications" in preparation, NBS Boulder, Colorado.

309
26 . Lackman, L. M. , Arvin, G. H. et al . , Advanced Composites Desig n Guide , 3rd Ed., Vol . II: Analysis, Air Force Materials Laboratory,

Wright-Patterson Air Force Base, OH (1973).

27. Wadoups, M. E. , "Characterization and design of composite materials",

in Composite Materials Workshop , Technomic Publishing Co., Inc.,

Standford, Conn. (1968) 254.

28. Loud, S. N. "Advanced composites applications of S-2 glass fiber".

Technical Paper, 34th Annual Technical Conference, Society of the

Plastics Industry, Section 20-F (1979) 1.

29. "Plastics for aerospace vehicles. Part 1: Reinforced plastics" MILHDBK-17A, Department of Defense, Washington DC (1971).

30. Lackman, L. M. , Arvin, G. H. et al . , "Advanced composite design

guide, 3rd edn. Volume IV: Materials", Air Force Materials Laboratory,

Wright-Patterson Air Force Base, Ohio (1973).

31. Handbook on materials for superconducting machinery" MCIC-HB-04,

Metals and Ceramics Information Center, Battel le, Columbus OH

(1977)

32. Hoggatt, J. T. "Development of cryogenic PRD-49-1 filament wound

tanks", NASA CR-120835, National Aeronautics and Space Administration,

Lewis Research Center, Cleveland, Ohio (1971).

33. Sims, D. F. and Halpin, J. C. "Methods for determining the elastic

and viscoelastic response of composite materials" ASTM STP 546,

American Society for Testing and Materials (1974) 46.

34. Softer, L. M. , Molho, R. "Cryogenic resins for glass filament-wound

composites", NASA CR-72114 (Final) National Aeronautics and Space

Administration. Lewis Research Center, Cleveland, Ohio (1967).

35. Kreibich, U. T., Lohse, F. , Schmid, R. "Polymers in low temperature

technology" Nonmetal 1 ic Material s and Composites at Low Temperatures " Plenum Press, NY (1979) 1.

36. Hartwig, G. "Mechanical and electrical low temperature properties

of high polymers" Nonmetal lie Materials and Composites at Low Temperatures " Plenum Press, NY (1979) 33.

37. Frank, W. F. X. "Spectroscopic analysis of high polymers for use at

low temperatures" Nonmetal lie Materials and Composites at Low

Temperatures Plenum Press, NY (1979) 51.

38. Schmid, R. "Correlation between valence energies and low temperature

flexibility" Nonmetal lie Materials and Composites at Low Temperatures

Plenum Press, NY (1979) 79.

39. Kneifel, B. "Fracture properties of epoxy resins at low temperatures"

Nonmetallic Materials and Composites at Low Temperatures Plenum Press, NY,

("1979) 123.

310
40 . Dahl rup-Peterson , K. "The ISR Cryogenic materials laboratory for

testing of composites, installation and results" Advances in

Cryogenic Engineering V 26 Plenum Press, NY (1980y!

41. Kasen, M. B., Schramm, R. E., Read, D. T. "Fatigue of composites at

cryogenic temperatures," ASTM STP 636, American Society for Testing

and Materials, (1 977) 141.

42. Must, J. G. , Arvidson, J. M. "Thermal conductivity of glass fiber/

epoxy composite support bands for cryogenic dewars" NBS report

275.03-78-2, National Bureau of Standards (1978).

43. Hust, J. G. , "Thermal conductivity" Semi-annual report on materials

research in support of superconducting machinery, NBS-ARPA report

NBSIR 76-848 (1976) 305.

44. Hust, J. G. , "Low-temperature thermal conductivity of two fiberepoxy composites" Cryogenics 15 (1975) 126.

311
TABLE I

ELECTRICAL PROPERTIES OF G-IOCR AND G-llCR

Temperature

(K)

Thickness

(cm)

Vol ume

Resi stivity

(n -cm)

Voltage Stress

at Breakdown

(kV/mm)

G -10 CR

295 0.030 48.4

0.036 8.9 X 10^^ 45.7

0.051 9.3 X

77 0.051 1.5 X

4 0.036 4.0 X 10^^ 48.4

0.051 4.1 X 10^^

G -10 CR

295 0.030 7.3 X

0.051 1 .3 X 10^^

77 0.051 6.1 X 10^^

4 0.030 2.0 X 10^^ 48.0

0.051 2.0 X 10^^

312
TABLE II

Typical Compositions of Glass Used as Reinforcement

for Cryogenic Laminates

(Weight Percent)

El ement E Glass Designation

S-910/S-2* R**

Si 02 54.4 65.0 60.0

AI

2

O

3

14.4 25.0 25.0

CaO 4.5 10.0 6.0

B

2 O3

8.0

Na 20 + K

2 O 0.5

•^^2^3 0.4

^2 0.3

*Owens/Corning Corp., Toledo, Ohio

**Veltroflex Div., Saint-Gobain Industries, Neuil ly-sur-Seine, France

313
ULTIMATE

TENSILE

STRENGTH,

MPa

X

h-O

liJ a:

cn

LlI

to2

liJ

UJ

»-<

2

3

Fig. 1 Temperature dependence of ultimate tensile strength of fabricreinforced polymer-matrix laminates. All data in warp direction

except for G-IOCR and G-llCR where fill direction strengths

are indicated by //// for 295 K and \\\\76/4 K. E-glass Fabric Bisphenol A epoxy

1 - (G-llCR), 67-72 w/o glass Ref. 6

2 - (G-IOCR), 64-68 w/o glass Ref. 6

3 - 74.8 w/o glass Ref. ^

4 - 62.5-66.1 w/o glass Ref. ^

E-glass Hat Bisphenol A epoxy

5 - Chopped-mat, 45-55 w/o glass Ref. 1J_

6 - Continuous Mat, 50-60 w/o glass Ref. 1

Miscellaneous E glass Fabric laminates

7 - polyimide, 66.9 w/o glass Ref. ^

3 - Melamine (G-5) Ref. 1J_

9

- Polyester Ref. 1

Other Laminates

10 - Cotton Fabric/Phenolic, (GE) Ref. 1J_ 11 - Pitch-Base Graphite Fabric/Bi

A epoxy, 61.8 w/o Graphite Ref. 9

12 - Aramid-Fabric/Bis A epoxy, 61.5

w/o A ram id Ref. 9

314
150

OIXLSd ‘H19N3yiS N0ISS3i;dW03 31VWI13n

ID

LD

CT) CM

OI

<o

s-OJ>

<x

"Oc

CO

r-^

OJ

cr.

<T3 S-<u>

ct

CO

1 ~

1

1

1

1

1

1

1 3

o

o

o

o

ID o

o

LD o

LD

CM

C|-

O

4-> C") c

OJ

J- t/l 4-> (U

(/) -M

(C c C

o •(

•r— E

(/5 fD

to 1

QJ

S_ X

C- • 1

E s_ o -M u (T3 E

cu 1 4-> s- (C QJ E E •1— >,

f

r— O

o c.

4- c

o cuu

OJ S- a o

c 4-

0> c

"O •1— c OJ

O) s_ O- 1 O) -M T3 03E

OJ

s_ -cc

+-> fO

fO

s_ u

OJ •r— Cl S- E

QJ 03 t—

CM

CD

CQ

U

j3

fO

,— 1

,— ,—

CO|CO “ "I

4- 4- 4- 4- 4- 4-

OI a' Q) QJ QJ QJ DC Dd DC DC DC d;

CO

CO

03 r—

CT) CO 00 CO 00 o

03 03 —

1— 1

5 LO —

CD CD X CO LlJ o LD —

C_) O o C- LD S- —

—

''— QJ 1 QJ QJ s 5 C -M o C •1

CO •1

CM CO E (U r—

LO 1

r- 03 o

1 1 o -t-J 1

t— c

c 03 QJ o cu LD LO OJ E E D.^

r

r« Q. •o 1/) t/l DC DC CO QJ CO tn 00 c O C_) • r— D. QJ fO rt3 o 1

O CD C. 4-> 1

14->

1

r— O 03 CD CD -M

1 1 d: C 1 1 o

CO CD tl5 C_3 •>- LlJ LU (_) z E

1 1 1 03 1 1 1 CO _J

1

C\J CO CO LO LO

fO S-QJ cr JC.

1 4-> LlJ O

edW ‘Hi9N3diS N0ISS3ddW00 3iVWIi3n

315
INTERLAMINAR

SHEAR

STRENGTH,

MPa

Fig. 3 Angular dependence of interlaminar shear strength of G-IOCR

glass-fabric epoxy laminate at 295 K and 76 K. Shaded areas reflect 2-sided, 95 % confidence intervals. Ref. 15 using

modified ASTM D 2733-70 method. NBS data for G-llCR indicated

by circles (6).

316
ra Q_

C£!

ZDQO

OO

</1 CL

QO

OO2: LU

Fig. 4 Temperature dependence of tensile moduli of fabric-and matreinforced polymer-matrix laminates. Waro direction of fabric

or edgewise direction of mat.

E-Glass Fabric Bisohenol A Epoxy

1 - G-llCR, 67-72 w/o glass Ref. 6

2 - G-IOCR, 64-68 w/o glass Ref. ^

3 - 74.8 w/o glass Ref. 9

4 - 62.5-66.1 w/o glass Ref. 9

E-Glass Mat Disphenol A epoxy

5 - Chopped Mat, 45-55 w/o glass Ref. Vl_

6 - Continuous Mat, 50-60 w/o glass Ref. 1

Miscellaneous E-glass Fabric Laminates

7 - Polyester Ref. Vl_

8 - Polyimide, 66.9 w/o glass Ref. £

Other Laminates

9 - Aramid/Epoxy 61.5 w/o fiber Ref. ^

10

- Pitch Base Graphite/Epoxy, 61.8 w/o fiber Ref. 9

317
Em

E(295)

Fig. 5 Temperature dependence of two principal Young's moduli for

G-llCR and G-IOCR glass fabric-reinforced epoxy laminates.

Resonance method (18).

318
TEMPERATURE, K

Fig. 6 Temperature dependence of modulus of the bisphenol A

resin systems used in G-IOCR and G-llCR glass-fabric

reinforced epoxy laminates. Resonance method (18).

epoxy

319
WARP

GPa

Fig. 7 Orientation dependence of room temperature Young's modulus

of G-IOCR gl ass-fabric/epoxy laminate. Resonance method (18).

320
- 0 .

- 0 .

- 0.8

1 2 3 4 5 6 7 8 9 10 11 12

Fig. 8 Temperature dependence of thermal contraction for fabric- and

mat-reinforced laminates. For fabric laminates, W = warp

direction, N = normal direction. For mat laminates, E = edgewise direction.

E-Glass Fabric Bisphenol A Epoxy

1 - G-llCR 67-72 w/o glass Ref. 6

2 - G-IOCR 64-68 w/o glass Ref. 6

3 - 64.4 w/o glass Ref. 9

4 - 66.1 w/o glass Ref. 9

5 - 67 w/o glass Ref. 9

6 - 62.5 w/o glass Ref. 9

7 - 74.8 w/o glass Ref. 9

E-Glass Mat Bisphenol A Epoxy

8 - Chopped flat, 45-55 w/o glass Ref. 11 9 - Continuous Mat, 50-60 w/o gla:5S Ref. 11 Mi seel 1 aneous Laminates

10 - Aramid/Epoxy , 61.5 w/o fi ber Ref. 9

11 - Pitch-base Graphite/Epoxy, 61 .8 w/o fiber Ref. 9

12 - Unreinforced Epoxy Ref. 9

321
THERMAL

CONDUCTIVITY

0.8 -

E

[liillll Approximate 295 K to 76 K Approximate 76 K to 4 K

8 8

Fig. 9 Temperature dependence of thermal conductivity for a series

of E-glass reinforced bisphenol A epoxy laminates and for

unreinforced bisphenol A epoxy. Data are illustrated for

fabric reinforcement (W = warp, N = normal) and for mat

reinforcement (E = edgewise, N = normal).

E-Glass Fabric Bisphenol A Epoxies

NBS Data

1 - G-llCR, 67-72 w/o glass Ref. 6_ 2 - G-IOCR, 64-68 w/o glass Ref. 5^ 3 - Bisphenol A epoxies Ref. 20

CERN Data (Ref. 9)

4 - 64.5 vj/o glass

5 - 62-74 w/o glass (5 laminates)

6 - 67.5 w/o glass

7 - Bisphenol A Epoxy

E Glass Mat Bisphenol A Epoxies

8 - Continuous-Mat, 50-60 w/o glass Ref. 11

322
0.8

liiil Approximate 295 K to 76 K

Fig. 10 Temperature dependence of thermal conductivity in the direction

normal to the reinforcement (flatwise direction) for a series

of silicon-, polyimide-, epoxy-, and ohenol ic-matrix laminates.

Ref.

E-Glass Fabric Laminates

1 - Silicone Matrix, 70 w/o glass

2 - Polyimide Matrix, 66.9 w/o glass

3 - Polyimide Matrix, 58.9 w/o glass

Miscellaneous Bisohenol A Laminates

4 - Aramid-Fabric , 61.5 w/o fiber

5 - Pitch-Base Graphite 61.8 w/o fiber

Phenolic Laminates 6-65 w/o cotton

323
TENSILE

STRENGTH,

MPa

TENSILE

STRENGTH,

MPa

no -

SOTO -

50 -

30 -

10 -

2000

o

0-

300,

200 -

40

30

ro2 295 K —Average 76 & 4K

20 a

TRANSVERSE

-1 16

12

ro

G1 ass Boron

H h

Graphite A ram id ten.

Epoxy

0

300

1500

1000

500 -

1 2 3

LONGITUDINAL

- 200

- 100

r

Fig. 11 Temperature dependence of the uniaxial longitudinal and

transverse tensile strengths of glass-, boron-, graphiteand aramid-reinforced epoxy laminates.

Glass Reinforcement

1 - S-901/Resin 2, 66 v/o fiber Ref. ^ 2 - R/Bisphenol A, 63 v/o fiber Ref. ^

3 - E/Bisphenol A, 50.6 v/o fiber Ref. £

Boron Reinforcement

4-5.6 mil B/2387 epoxy, 52 v/o fiber Ref. ^ 5 - 5.6 mil B/6061 aluminum, 47 v/o fiber Ref. 23

Graphite Reinforcement

6 - AS/Resin 2, 64 v/o fiber. Ref. ^ 7 - HM-S/934, 62 v/o fiber Ref. ^

8 - HM-S/Bi sphenol A, 55 v/o fiber Ref. ^

9 - GY-70/934, 65 v/o fiber Ref. ^

Aramid Reinforcement

10

- Kevlar 49/934, 67 v/o fiber Ref. 2^

324

TENSILE

STRENGTH,

10

psi

TENSILE

STRENGTH,

10
ULTIMATE

TENSILE

STRENGTH,

MPa

200

AS Fiber (hiqh strength)

1200

1000

800

600

400

HM-S (medium strength & modulus)

TEMPERATURE, K

Fig. 12 Comparison between the

temperature dependence

of graphite-reinforced

obtained at NBS (23,25

combined literature data (2) on the

of uniaxial ultimate tensile strength

laminates (shaded area) and the data

) and CERN (9).

325

ULTIMATE

TENSILE

STRENGTH,

10"

psi
100

o

Q_

X

h-o

2

UJtr

h—(O

(/) (O

UJX

2

o

o

_l

<

X

<

X

<

Glass Boron Graphite Ararrnd

Fig. 13 Temperature dependence of the uniaxial longitudinal and

transverse compressive strengths of glass-, boron-, graphiteand ararrid- reinforced laminates. Plotted values are maximum

values for the longitudinal direction and average values for

the transverse direction.

Glass Reinforcement

1 - S-901/Resin 2, 66 v/o fiber Ref. 23

Boron Reinforcement

2 - 5.6 mil B/2387 epoxy, 52 v/o fiber Ref. 2^

3 - 5.6 mil B/6061 aluminum, 47 v/o fiber Ref. 23

Graphite Reinforcement

4 - AS/Resin 2, 64 v/o fiber Ref. ^ 5 - HM-S/934, 62 v/o fiber Ref. n

6 - GY-70/934, 65 v/o fiber Ref. ^

Aramid Reinforcement

7 - Kevlar 49/934, 67 v/o fiber Ref. 25

326

MAXIMUM

UNIAXIAL

COMPRESSION

STRENGTH,

10
TENSILE

MODULUS,

GPa

TENSILE

MODULUS,

GPa

140

10

0

295 K Average 76 & 4 K

TRANSVERSE

w

Q.

too

CO3

_Jz>QO

UJ

to2

UJ

h1 1 I 1 I 1

I 1 I 1 Bis A

Epoxy

50

40 i;; Q.

COo—

30 to" 3

3

QO

- 20

UJ

CO ^ UJ 10 I0

Figure 14 Temperature dependence of uniaxial longitudinal and transverse

tensile moduli of glass- boron- graphite- and aramid-reinforced

laminates.

Glass Reinforcement

1 - S-901 /Resin 2, 66 v/o fi Ref.

^

2 - R/Bisphenol A, 63 v/o fiber Ref. ^

3 - E/Bisphenol A, 50.6 v/o fiber Ref. ^

Boron Reinforcement

4 - 5.6 mil B/2387 epoxy, 52 v/o fiber Ref. ^

5 - 5.6 mil B/6061 aluminum, 47 v/o fiber Ref. ^

Graphite Reinforcement

6 - AS/Resin 2, 64 v/o fiber. Ref. n

7 - HM-S/934, 62 v/o fiber Ref. 2S

8 - HM-S/Bisphenol A, 55 v/o fiber Ref. 9

9 - GY-70/934, 65 v/o fiber Ref. ^

Aramid Reinforcement

10

- Kevlar 49.934, 67 v/o fiber Ref. 25

327
IN-PLANE

SHEAR

MODULUS,

GPa

12.0

IQ. 0

8.0

6.0

4.0

2.0

0

U)a. o

oo

QO

Od

oo

<

I Cl.

Fig. 15 Temoerature dependence of in-plane shear modulus, 0,^, of

uniaxial glass-, boron-, graphite- and aramid-reinforced epoxy

laminates (23).

1 - S-901-glass/Resin 2, 66 v/o fiber Ref. ^ 2 - 5.6 mil 8oron/2387 epoxy, 52 v/o fiber Ref. ^

3 - AS Graphite/Resin 2, 64 v/o fiber Ref. ^

4 - HM-S graphite/934 epoxy, 62 v/o Ref. ^

5 - GY-70 graphite/934 epoxy, 65 v/o Ref. ^

6 - Kevlar 49 aramid/934 epoxy 67 v/o Ref. ^

328
Epoxy

Fig. 16 Temperature dependence of thermal expansion of uniaxial glass-,

boron-, and graphite-reinforced laminates in the longitudinal

(L) and transverse (T) directions. Epoxy matrices are a con- ventional bisphenol A (Bis A) or a flexibilized formulation

(Resi n 2 , Ref. . Glass Reinforcement

1 - S-901/Resin 2, 66 v/o fiber Ref. ^ 2 - R/Bisphenol A, 63 v/o fiber Ref. ^

3 - E/Bisphenol A, 50.6 v/o fiber Ref. ^

Boron Reinforcement

4 - 5.6 mil B/2387 epoxy, 52 v/o fiber Ref. ^

5 - 5.6 mil B/6061 aluminum 47 v/o fiber Ref. 23

Graphite Reinforcement

6 - AS/Resin 2, 64 v/o fiber. Ref. ^

7 - HM-S/Bi sphenol A, 55 v/o fiber Ref. £

Epoxy Resins

8 - Bisphenol A Ref. ^

9 - Resin 2 Ref. 20

329
THERMAL

CONDUCTIVITY,

W/mFig. 17 Temperature dependence of thermal conductivity of uniaxial

laminates. L = Longitudinal, T = Transverse. Transverse

conductivity of the aramid laminate assumed from fabricreinforced data.

1 - S-glass/epoxy , 65 v/o fiber Ref. ^ 2 - S-glass/epoxy, 63 v/o fiber Ref. ^

3-4 -5.6 mil Boron/epoxy, 52 v/o fiber Ref.

5-6- 5.6 mil Boron/aluminum, 47 v/o fiber

7 - Graphi te/epoxy Ref. ^

8 - Graphite/epoxy Ref. ^

9 - Aramid/epoxy Ref. ^

10 - Aramid/epoxy Ref. ^

11 - Bisphenol A epoxy Ref. ^

Ref. 43

330
UNIAXIAL

COMPRESSION

STRENGTH

CHANGE,

MPa

250

200

150

100

50

0

-50

Fig. 18 Range observed in uniaxial longitudinal compression strength

at 295 K, 76 K and 4 K for a series of laminates. All data are normalized at 295 K.

1 - HM-S Graphite/934 epoxy

2 - GY-70 Graphite/934 epoxy

3 - AS Graphite/Res 2 epoxy

4 - S-901 Glass/Res 2 epoxy

5 - 5.6-mil Boron/6061 aluminum

6 - 5.6 mil Boron/2337 epoxy

Ref. 25

Ref. 25

Ref. 23

Ref. 23

Ref. 23

Ref. 23

331

UNIAXIAL

COMPRESSION

STRENGTH

CHANGE,

psix

10^

MECHANICAL PROPERTIES OF G-IOCR AND G-llCR

National Bureau of Standards

333

Mechanical Properties of G-IOCR and G-llCR

R. E. Schramm and M. B. Kasen

The last report contained mechanical characterization data at 295,

76, and 4 K in the warp and normal directions for G-IOCR and G-llCR from

manufacturer A's pilot plant. This characterization has now been

completed for the fill direction. The results of all the individual

tests are in Table 1. Three additional batches of material have been

screened to examine variations between batches and manufacturers:

1) manufacturer A, production plant, 2) manufacturer B, production

plant, and 3) manufacturer B, boron-free glass (a special variant which

might display less susceptibility to radiation damage). This testing

covered only two specimens in each of the configurations at only 295 and

76 K. These results are listed in Table 2.

All of the mechanical data were summarized with most of the electrical

and thermal results in a paper presented at the International Cryogenic

Materials Conference at Madison, Wisconsin in August, 1979. This paper

will be published in Advances in Cryogenic Engineering, Vol . 26 . This

paper discussed our findings and a preprint follows.

As part of a program for the USAE Aero Propulsion Laboratory, we

examined the effects of thermal ageing. Short-beam shear specimens of

G-IOCR and G-llCR from Manufacturer B were heated to 122°C for 2 h and

112 h and then tested along with untreated specimens at 76 K. The

results in Figure 1 indicate no effect on G-IOCR. The strength drop

noted in G-llCR after 112 h might indicate some deterioration, but this

may not be statistically significant given the normal data scatter.
CO

COQ.

COO

I

I-O

LU

CCH

W

DC<

111 X

w

<

LUm

CO0

1

10

Figure 1. Effect of aging on short-beam shear strength.

In cooperation with MIT, we will conduct interlaminar shear tests on

1.3-cm (0.5-in) thick G-IOCR guillotine specimens at 4 K and then compare these results from 0.38-cm (0.15-in) thick material. Additional

work will include fatigue cycling of both shear and compression specimens at 4 K and measurement of residual strength.

Associated Programs

The following NBS activities in the nonmetallics area are of potential

interest to MFE cryogenic technology:

1. Characterization of the cryogenic properties of high-modulus,

advanced composites and the assessment of their applicability to lightweight MHD and rotating machinery superconducting magnets (USAF Aero

Propulsion Laboratory).

336
2. Interaction with a leading industrial glass research laboratory to disseminate information on the cryogenic mechanical thermal

properties of glass mat-reinforced epoxy laminates developed and characterized by this laboratory. These data will appear in the second

supplement of the LNG Users Manual.

3. Evaluation of the various mechanical properties and coefficients

of friction of polymeric materials vs. metals at cryogenic temperatures

(Air Force Materials Laboratory).

FY 80 Nonmetal lies Program

The FY80 nonmetal lies program will continue to reflect the needs of the

MFE program as defined by ongoing program interaction and suggested by

input from the Vail NBS/DoE workshop. Three major areas are presently

defined:

1. Increasing the understanding of the fundamental relationship

between cryogenic temperature and the performance of nonmetal lie composites.

2. Increasing the understanding of the effect of neutron and

gamma radiation on the performance of nonmetal lie composites at cryogenic

temperatures.

3. Continued interfacing between manufacturers of cryogenic

composite laminates and the fabricators of MFE superconducting magnet

systems.

337
INITIAL

MECHANICAL

PROPERTY

VALUES

OF

G-IOCR

AND

G-llCR

(MANUFACTURER

API

LOT

PLANT)

^ ^ ^

r— voIld cm 00

V£) CO LT> ^ a^

[CM CM CM CvjICVJ

^l^o t CM CM CM CM CM CM CM CM

COO, O I— •— I— CO 00 100 00 00 <

I— *£>

I VO 0^ CM 0^ «

COQ_

, CM O ^ O O I—ko o ^ o VO ^ VO kr> • • • .

. . . . CM 'd- CO

c£) CO volvo I— I— »— I—

CO Vf) ^ IT) I

jlvn VO VO '0|VO >|vo CO CO CO CO CO CM CO I

s- o

to x:

0) CO

VO VO VO CM LT) VO VO ^ CO CM CO < o ^ r**. O ^ VO o> <

oooLOi— vncMco 00 voor^oo

COOQO

p^ cr. CM

C_D ro CO CM CM

CO P««* CO VO P^ VO

CO »— CO CO

00 VO p^ VO

CM ID ir> LO

I— VOCOP^ P^ CM VO CM U0OVCr*v

CO

p^

VO00 O VO

00 ^^ r- O

P«N. CO ^ < O CTi I vn VO VO I

LO^COP^ ^OCOCM r^OLO^

CM P^ CO ^ VO P^

VO lO VO ^ VO 00 CO < OS VO ^ 00 < VO VO VO VO <

CM CO CO O CO VO VO 00 CO P«** OS P*^ I— ^ OS OS ^ I— 0> p***. CM «

(/) OCO O p^ t VO VO I CM CM CM «

If— P^ VO OS CO VO CM

VO •— OS ^ ^ ^

VO CM O ' OS I— 00 < ^ VO ^ •

Si— O OO 00 ^ VO < CO CM CM < CO CO CO <VO CM 00

VO f— •— < O p**.. VO t VO VO VO I

00 O VO OS » CO VO ^ « VO VO I lO VO VO I

- >

co

' o o o o

^ ^ CM ' VO VO I— . •— CM CM ' ^ ^ VO * O CM I— I CM CM CM «

VO 00 CM I I— O CM • CM CM CM <

I o o o o <

COo.

CO ^ CO CO |co «?j- cj- CO CO CO CO

CM VO VO LO VO I CO CO P*«* OS

338

(F)

=

Fill

direction

of

reinforcement

*

=

Calculated

values
1^ CO ro|ro r— f— ], • |r— CO ro|ro |r— CO foiro , •|i— fO COlfO

I O

o

o

c

o

COQ. «>olr^ *— o^lo VO VOlvD 00 VO CTi oolco ID inlLO COO'IOO vo vOlvO 0> ^ I ir> voim 00 oolco

cc: o

• CM r«H CO Ko «3- o CM CO CO CO o oUr> o> col^ r— CO ^ CM CO •— kj- r-s. O

«:r *>o r^iP^

*— Kr <— *— •—

Q

2: <c

cxl o

o

I C3

COCL

o in CO

CO oolco

r>. r«* CM lo I

o^|Q^ in VO »

•5J- CM < a^ CM <

00 colco CO C7\

cr> VO CO ^ o.rv.

vn VO m CO co « O' O'lo' r*** r>. t 00 oolco O' CT' < O CO CM CM I— t

CO O'ico vO VO * ^ O CM O O' I

(U rH

.od

UJo

00OCt^

o> rj X

> o •»- D‘«^ 5 00

<U CT> Q_

L C

Q. OJ S S- ro o r\

u CO CM CO lO O CO

I o vn m o <c o ' VO CO m CO O O'

VO VO ^ ID ^ i CD VO CO I i- CO ^ <

) vn CO CM o I— O' O' O' CO O' CO co <— m

CO CM

VO VO

I

VO lO

CO v-> O' <

. ro • •

I— 4- CO CO f CO 3 CM »— < VO C <3- «

[i— S- O <— f—

CM CO O CO CO

VO r«-.Ir^

»- CO CM < rj VO CO • c «?} »sr «

3 CO CO vO

I— m ( P^ O' < Ol CO CO <

l_ CO CO CO c**. CO *

> r- o VO CO cm*—'

O' O' O'

in CO ^ p^ p^

s: o lo < ' VO CM • CO CO <

CM in < in o < CO 00 (

S CM CM CM ^ < •lo CO cmIco CO in co cm

</) Q.

r- r- ^ o

a:UJ

Q_OCiT QCO or*^vOr^ CMOvo

> in CO r— i

I I— CO <

• »— CM CM «

CM CM < in p^ « O CO f CM < O' O' (

P*^ CM

t— CO

CM CM

px. r— O' CM IP-- I— CM f— r-v co|r-s

<3- CO colco

CM p— f— t

r- O f— CM CM CM O O

CO oolio o olin

O' r-H CO m o ^ c^> o cr»

CM r— .

P-. CO ‘ CM CM t CM CM I

P-. O' 00

in in in CO colco O CO rin vo m

VO volvo

CO 'O- CO CO CM CM

in CO r-. vD 00 r-

<3- «3-I^

CM r—

CO 00 ( o o <

«3‘ O' CM

<3‘ *—^00

O' © O'

lO O < CM «3- t CM CM (

</) O.

rO «3’|<3’ ^ ^1 ^ ^ ^l«3“ m ^|m «3‘1^ <3- ^ ^1^ in inlin fO colco ^ ^|«^ ^ injin in mlm

nr CO O' nr CO nr in O' CM CO f— P-- CO CM m O CM ^ O' ^ O 00 O' O' O O' O O — lO

339

(W)

=

Warp

direction

of

reinforcement

*

=

Calculated

values

**

=

Broke

in

tension

-

not

shear
:

(
MECHANICAL, ELECTRICAL, AND THERMAL CHARACTERIZATION OF

G-IOCR AND G-llCR GLASS CLOTH/EPOXY LAMINATES

BETWEEN ROOM TEMPERATURE AND 4 K

National Bureau of Standards

341

MECHANICAL, ELECTRICAL, AND THERMAL CHARACTERIZATION OF

G-IOCR AND G-llCR GLASS CLOTH/EPOXY LAMINATES

BETWEEN ROOM TEMPERATURE AND 4 K*

M. B. Kasen, G. R. MacDonald, D. H. Beekman, and R. E. Schramm

Fracture and Deformation Division

National Bureau of Standards

Boulder, CO 80303

ABSTRACT

The magnetic fusion energy program of the U.S. Department of Energy

will require large-scale use of industrial high-pressure fiberglass/

epoxy laminates for structural support and electrical and thermal insulation.

In response to the need for commercially available and predictable

materials, the National Bureau of Standards and industrial suppliers

have cooperated with the National Electrical Manufacturers' Association

(NEMA) to establish specifications for cryogenic grades of G-10 and

G-11 glass-cloth laminates. These materials, designated G-IOCR and

G-llCR, differ from conventional laminates in that they are produced to

the same material and process specifications by all commercial laminators. Both products meet existing NEMA, federal, and miltary specifications. We have characterized the 295, 76, and 4 K mechanical properties of G-IOCR and G-llCR produced by two manufacturers. Properties

studied include tensile strength, elongation, modulus, Poisson's ratio,

and compressive strength. Interlaminar shear strength was characterized

by both guillotine and short beam methods. The effect of orientation

was examined. The data indicate that both laminate types are excellent

cryogenic materials. Additionally, we have screened the mechanical

properties of the same laminates produced with boron-free E glass reinforcement in anticipation of improved radiation resistance of this

variant.

*Supported by the Office of Fusion Energy, U.S. Department of Energy.

343
INTRODUCTION

The U.S. magnetic fusion energy program will require large-scale

use of industrial high-pressure glass-reinforced epoxy laminates for

structural supports and for electrical and thermal insulation in superconducting magnets. This choice is based on such factors as availability,

cost, machinabil i ty, high strength-to-weight ratio, and low electrical

and thermal conductivities. There is an immediate need for a commercial

supply of such materials having predictable cryogenic performance. NBS,

therefore, has cooperated with the U.S. laminating industry and the

National Electrical Manufacturers' Association (NEMA) to establish

specifications for cryogenic grades of NEMA G-10 and G-11 glass cloth

reinforced epoxy laminates. These materials, designated G-IOCR and GIICR, meet existing NEMA, as well as federal and military, specifications for G-10 and G-11 products. The component and manufacturing

specifications have been distributed to all U.S. manufacturers through

the NEMA organization. The products are currently available from one

manufacturer, and others have shown an interest in their production.

NBS is currently undertaking an electrical, thermal, and mechanical

property characterization of the CR-grade materials over the 295 to 4 K

temperature range. Contributions to the characterization have also been

made by universities, national laboratories, and a private corporation.

In this paper, we present the data base obtained on G-IOCR and G-llCR

produced for NBS by two manufacturers. For comparison, we screened the

mechanical properties of CR-grade laminates produced with boron-free

glass reinforcement in anticipation of improved radiation resistance of

this variant.

344
MATERIALS

The basis for both G-IOCR and G-llCR is the same glass fabric of

silane-finished E glass. Its plain weave is produced on a conventional

loom by interlacing warp (length) threads (16. 9 ± 1 .2 per cm or 43 ± 3

per inch) and fill (width) threads (12.6 ± 0.8 per cm or 32 ± 2 per

inch). The specially made boron-free E glass is similarly configured.

The G-IOCR matrix is a conventional heat-activated, amine-catalyzed

bisphenol A solid-type epoxy resin chosen for proven performance at

cryogenic temperature. The G-llCR matrix is an aromatic-amine hardened

bisphenol A liquid-type epoxy resin expected to provide improved resistance to radiation damage. The finished products are a natural beige

color. Resin weight fraction is 32 to 36 % for G-IOCR and 28 to 33% for

G-llCR. Production details are presented in another paper [1].

Two producers, designated manufacturers A and B, supplied sheets of

G-IOCR and G-llCR in nominal thicknesses of 0.127 cm (0.050 in), 0.0381 cm

(0.150 in) and 2.54 cm (1 in). Additionally, manufacturer B provided

similar panels in the boron-free variant. Manufacturer A produced the

materials as a pilot plant run. Although pilot plant procedures closely

parallel production conditions, the manufacturer subsequently provided

additional actual production panels for additonal testing.

The mechanical tests performed on these materials are summarized in

Tables I and II. As noted on Table I and Figures 1 through 3, the pilot

plant production of manufacturer A was used to establish base mechanical

properties for tension (warp, fill), compression (warp, fill, normal)

and shear (warp, fill). Results of the screening tests on the mechanical performance of material from manufacturer B appear in fable II.

345

50

G-10 CR

H Warp

Fill

G-llCR

J Warp

Fill

Figure 1. Young's modulus versus temperature for G-IOCR and G-llCR

(Manufacturer A, Pilot Plant). Error bands show uoper and lower

data spreads.

347

YOUNG'S

MODULUS

Figure 2. Tensile strength versus temperature for G-IOCR and G-llCR

(Manufacturer A, Pilot Plant). Error bands show upper and lower

data spreads.

349

COMPRESSIVE

STRENGTH,

\ Q.

CL COO

o

Figure 3. Compressive strength versus temperature for G-IOCR (Manufacturer A, Pilot Plant). Error bands show upper and lower

data spreads. G-llCR displayed similar data spreads.

351

COMPRESSIVE

STRENGTH,
I

I
These tests were performed only in the warp direction. The pilot plant

production of manufacturer A was also used in the study of electrical

performance (Table III), thermal conductivity (Figure 4), and thermal

expansion ( Fi gure 5)

METHODS

MECHANICAL TESTS

Tensile specimens cut from the 0.127-cm sheets were a modified

version of the MIL-HDBK 17A specification. The gage length was 5.4-cm

(2,125-in) long and 1 .27-cm (0.5-in) wide. The specimen width tapered

to 2.54 cm (1 in) in the grip area. Resistance strain gages cemented

directly to the surface in longitudinal and transverse orientations

allowed measurement of Young's modulus and Poisson's ratio. The ultimate strength and elongation were also determined.

Warp and fill compression strength specimens from the 0.381 cm

sheets were 6.99-cm (2.75-in) long and 0.794-cm (0.313-in) wide; specimens tested normal to the glass cloth had about the same cross-section

but were only 2.54-cm (1-in) long. With end caps cemented in place, the

gage length was 1.91-cm (0.75-in). The compression method and fixture

were described elsewhere [2].

Interlaminar shear strength tests followed two methods. Threepoint bending (ASTM D 2344-72) was the simplest and good for comparison

or quality checks but provided no real engineering data, since tensile

and compressive stresses were also present. The specimens were 2.54-

cm (1 in) long by 0.64-cm (0.25-in) wide and came from the 0.381 cm

thick sheets. The guillotine shear method (ASTM D 2733-70) can provide

more meaningful values of strength, but it, too, is subject to considerable inaccuracy unless care is used in its application. Data appearing

353
1000

i.)l I.IU MU* 'AilAliOnONOO 1VI/\IU3Hi

354

Figure

4.

Thermal

conductivity

versus

temperature

for

G-IOCR

and

fl-lTCR

(Manufacturer

A,

Pilot

Plant).
THERMAL

CONTRACTION

(10“AL/L)

Figure 5. Thermal contraction of G-IOCR and G-11CR. The discrete points

were from the preliminary tests done at LASL while the continuous

curves were measured at NBS. (Manufacturer A, Pilot Plant).

355
on Tables I and II were obtained with a modified ASTM specimen (saw cuts

on opposite surfaces were 0.953 cm (0.375 in) apart with side support

established at a level yielding consistent results regardless of specimen

thickness. Without this support, an apparent peeling of each side from

the shear plane reduced the measured strength by 50%.

Conditioning of all specimens prior to testing (ASTM D 3039-71)

was ^ 40 h at 22 ± 2°C and 50 ± 10% relative humidity.

ELECTRICAL TESTS

Volume resistivity tests conducted by the subcontractor were in

accord with ASTM D 257 with 1000 Vdc and 7.62 cm (3 in) diameter electrodes with a 0.64-cm (0.25 in) guard gap. For the electric strength,

60 Hz was applied at 0.5 kV/s rise with 1 .27-cm x 3.81 cm (0.5-in x

1 . 5-pai nted el ectrodes

THERMAL TESTS

Thermal conductivity measurements were conducted by the longitudinal heat flow method at closely spaced temperature [3]. The spec3 imens were 1.91 cm (0.75 in) cubes with a density of 1.904 g/cm for

G-IOCR and 1.956 g/cm^ for G-llCR.

A preliminary test of total contraction between room temperature

and 75 K was conducted at the Los Alamos Scientific Laboratory. Later

measurements as a continuous function of temperature were performed at

NBS with a quartz dilatometer.

RESULTS AND DISCUSSION

MECHANICAL TESTS

The mechanical data in Table 1 are a characterization of the pilot

plant G-IOCR and G-llCR from manufacturer A at all three temperatures

and in both fiber directions (also normal direction for compressive

356
strength). These are averages of three or four specimens. The standard

deviations among individual tests were about 4% of the average value for

the Young's modulus, tensile strength, and shear strengths, and about 9%

for the Poisson's ratio, compressive strength, and strain-to-fai 1 ure

Overall, there was no noticeable change in data scatter with temperature. In Figures 1 through 3, the shaded areas indicate the extent of

upper and lower data spread among tests.

Only two parameters showed any significant temperature related

material differences. G-llCR had a higher modulus in both directions

and a higher tensile strength in the fill direction. The guillotine

shear strength may indicate some superiority in the G-IOCR.

Cooling both materials from room to cryogenic temperatures increased the value of all parameters: Young's modulus about 30%, Poisson's

ratio 35 to 40%, tensile and compressive strengths 70 through 130%,

ultimate strain about 100% in the warp direction and 60% in the fill

direction, short-beam shear strength about 100%, and guillotine shear

strength in the range of 50%. G-llCR appears to be slightly less temperature dependent than G-IOCR. The elastic properties, modulus, and

Poisson's ratio showed somewhat greater change with temperature for fill

than warp directions, while the opposite occurred for the ultimate

strength and elongation.

Since the fill direction yarn count is approximately 0.75 of the

warp direction count, it would seem likely to find the fiber-dependent

parameters occurring in roughly this ratio. Fill/warp was about 0.80

for the modulus and ultimate strain, 0.65 for tensile and compressive

strengths, and 1.00 for Poisson's ratio. Shear strength is largely a

function of matrix and matrix-fiber bonding, so it should be directionally

357
00 LD 0 »— to u» ‘t- a. CM 00 LD r's. 0> u- 5C C'K LD UO c c

o»U 4-> UO 00 *— D r— 00 CO UD LD LD CM

L.(D 3 0. CVJ r— CM 0 LD LD

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x:

—

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m «=3- CM <n

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H-Z

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51 0

CO 0 0 ro. CO CM

CO CL UO CTt LD CM

•*— S- CM CM CM 0 to 0. 3 CD 0 0 0 0 0

LU «=r 0 UO cn — 03

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358
independent. The guillotine tests showed no real difference between

orientations, but the short-beam tests yielded a fill/warp property

ratio of 0.70; this is indicative of flexure-producing tensile and

compressive as well as shear forces. Some material dependence with the

fill/warp ratio did appear in the tensile and compressive strengths,

where it was somewhat higher for G-llCR than for G-IOCR. The only

parameter to display a temperature dependence was the failure strain,

the ratio of fill/warp properties decreasing in the cryogenic region.

The compressive strengths normal to the cloth plane are valuable

data for high-pressure laminates, since these materials are likely to be

thus loaded in many weight-bearing applications. In general, the values

seem to be comparable to those in the warp direction and show the same

temperature dependence.

The screening data in Table II include only two temperatures and

only the warp direction; they do, however, cover composite lots from

manufacturer A and B's production plants as well as the boron-free

product from manufacturer B.

Some differences exist between values in the two tables. All

comparisons here are to the complete characterization data in Table I

for manufacturer A--pilot plant. For G-IOCR, manufacturer A--production

plant, Poisson's ratio runs about 15% high while the compressive strength

at 76 K is 15% low. For G-llCR, manufacturer B--production plant, the

tensile strength at 76 K is 18% high.

The majority of differences, however, occur in the boron-free

products. In G-IOCR the modulus is 15% low, Poisson's ratio is 23%

high, tensile strength is 14-20% low, and failure strain at 76 K is 17%

low. The differences in G-llCR are a 9% high modulus, and the 76 K

359
TABLE

II

MECHANICAL

SCREENING

OF

G-IOCR

AND

G-llCR

THREE

BATCHES

(AVERAGE

OF

TWO

SPECIMENS)

COo

ol

Q. ro

i- Q.

ro O

360

(1

MPa

=

145

psi)

(1

GPa

=1

.

45x1

0^

psi

)
tensile strength is 19% high. None of these differences are really

dramatic, and they could be the result of normal experimental scatter

and the allowable variations in yarn count and resin content.

Limited data, particularly at cryogenic temperatures, are available

for comparison, and they deal mostly with S glass cloth [4,5]. In

general, they show that the present materials display a higher modulus

and tensile strength by a factor of about 1/3 and a somewhat higher

compressive strength and lower’ fail ure strain. Direct comparison of

shear strengths is very difficult because of the many experimental

variabl es

ELECTRICAL TESTS

Volume resistivity (Table III) showed 1.5 to 2 orders of magnitude

increase on cooling to 4 K. Electric strength, however, displayed no

temperature dependence and no difference between composite types.

THERMAL TESTS

Thermal conductivities (Fig. 4) were nearly identical for both

materials. Data were about 10% higher for the warp direction than normal

to the reinforcement. Fill direction measurements would likely yield

values similar to present results. Uncertainties in measurement are

about ± 5%. Complete details and analysis will be published later [6].

The thermal contraction data (Fig. 5) show a directional dependence

possibly indicating that the epoxy was a dominating factor. A detailed

analysis of the continuous data is now in preparation for publication

elsewhere [7]. A preliminary measurement of total contraction between

room temperature and 75 K conducted at LASL indicated approximately 10%

greater aL/L.

361
TABLE III

ELECTRICAL PROPERTIES OF G-IOCR AND G-llCR

MANUFACTURER A-PILOT PLANT

Temperature

(K)

Thickness

(cm)

Volume

Resistivity

($7 -cm)

Voltage Stress

at Breakdown

(kV/mm)

G-10 CR

295 0.030 48.4

0.036 8.9 X 10^^ 45.7

0.051 9.3 X 10^^

77 0.051 1.5 X 10^^

4 0.036 4.0 X 10^^ 48.4

0.051 4.1 X 10^^

G-11 CR

295 0.030

1 5 7.3 X lo'^

0.051 1.3 X 10^^

77 0.051 6.1 X 10^^

4 0.030 2.0 X 10^^ 48.0

0.051 2.0 X 10^^

362
CONCLUSIONS

The CR grade glass-epoxy laminates should prove useful in those

applications that require a structural or insulating material of predictable performance that is commercially available. These composites

do have good mechanical properties at room temperature and below. In

this and other laboratories , work continues on mechanical, electrical,

and thermal characterization as well as resistance to radiation damage.

’ACKNOWLEDGMENT

We are grateful to the Office of Fusion Energy, U.S. Department of

Energy for sponsoring this work. Members of the NBS Thermophysical

Properties Division performed the thermal conductivity studies (J. G.

Must) and thermal contraction measurements (A. F. Clark and M. A. Ranney).

We also wish to thank G. A. Mullen of the Westinghouse R&D Center,

Pittsburgh, for the electrical data and J. Harlow and R. I. Schermer of

the Los Alamos Scientific Laboratory for the preliminary thermal contraction data.

363
nl

, u^'hMU' ^i0i^i^m,

• y ., • (iK| -• •:•<*. ‘ ?'<:'

^*St

*l»<lll m

^'vT'

r4?i

( %‘^n(\ <s5 ^'

. 4 ^ ^fim

m

^ i S‘HU}^ 1^%,' '-mi

•r.m

./tpeiHW' -»^v . , *'' - .

i

•Y,v^- ;-i '.

;

v¥\^ ' M

''.V'''r ^''' r\. \-y y-’ , ••/li’.'^'v.
REFERENCES

1. J. R. Benzinger, in Advances in Cryogenic Engineering -- Material

,

Vol . 26 Plenum Press, New York (1980), in press.

2. R. E. Schramm and M. B. Kasen, Mat . Sci . Eng . 30:197 (1977).

3. J. G. Must and J. M. Arvidson, "Thermal Conductivity of Glass

Fiber/Epoxy Composite Support Bands for Cryogenic Dewars," Unpublished

report 275.03-78-2, National Bureau of Standards, Boulder, Colorado.

4. M. B. Kasen, Cryogenics 15:327 (1975).

5. M. B. Kasen, in Composite Reliability , ASTM STP 580, American

Society for Testing and Materials, Philadelphia, PA (1975), p. 586.

6. J. G. Must, private communication, NBS, Thermophysical Properties

Division Boulder, Colo. 80303.

7. A. F. Clark, private communication, NBS, Thermophysical Properties

Division, Boulder, Colo. 80303.

365
LIST OF TABLES

Table I.

Table II

Table II

Figure 1

Figure 2

Figure 3

Figure 4

Figure 5

Mechanical Characterization of G-IOCR and G-llCR, Manufacturer APilot Plant (Average of at Least Three Specimens).

Mechanical Screening of G-IOCR and G-llCR, Three Batches

(Average of Two Specimens).

Electrical Properties of G-IOCR and G-llCR, Manufacturer A- Pilot

PI ant.

LIST OF FIGURES

Young's modulus versus temperature for G-IOCR and G-llCR

(Manufacturer A, Pilot Plant). Error bands show upper and lower

data spreads.

Tensile strength versus temperature for G-IOCR and G-llCR

(Manufacturer A, Pilot Plant). Error bands show upper and lower

data spreads.

Compressive strength versus temperature for G-IOCR (Manufacturer A, Pilot Plant). Error bands show upper and lower

data spreads. G-llCR displayed similar data spreads.

Thermal conductivity versus temperature for G-IOCR and G-llCR

(Manufacturer A, Pilot Plant).

Thermal contraction of G-IOCR and G-llCR. The discrete points

were from the preliminary tests done at LASL while the continuous

curves were measured at NBS. (Manufacturer A, Pilot Plant).

366
TEMPERATURE DEPENDENCE OF YOUNG'S MODULUS AND

INTERNAL FRICTION OF G-IOCR AND G-llCR EPOXY RESINS

National Bureau of Standards

367

TEMPERATURE DEPENDENCE OF YOUNG'S MODULUS AND INTERNAL FRICTION

OF G-IOCR AND G-llCR EPOXY RESINS*

H. M. Ledbetter and G. Maerz

Fracture and Deformation Division, NBS, Boulder, Colorado 80303

Abstract

The Young's moduli of the epoxy-resin matrix material used in NEMAdesignation G-IOCR and G-llCR fiberglass-cloth-reinforced composites

were measured dynamically and semiconti nuousl y between ambient and

liquid-nitrogen temperatures. Both materials exhibit regular temperature behavior, showing large Young ' s-modul us changes, about 125 and 50

percent, respecti vely. Internal friction decreased about 80 percent

during cooling to liquid-nitrogen temperature (76 K) . The different

thermoelastic coefficients of the two materials indicate a different

internal structure.

Key words: Composites; elastic constants; epoxy resin; internal friction;

low temperatures ; Young's modulus.

369
Introduction

Large quantities of low-cost, high-pressure industrial laminates

are now used as structural members and as electrical insulators. The

National Electrical Manufacturers Association (NEMA) has established

composite-material grades G-10 and G-11 for these purposes. These

materials consist of glass-cloth-fiber reinforcement in an epoxy-resin

matrix. Cryogenic grades of these materials are designated G-IOCR and

G-llCR.

Previously, Ledbetter^ summarized the importance of elastic-constant

information for composite materials; and he reported Young's modulus and

internal friction for both G lOCR and G-llCR composites between room

temperature and liquid-nitrogen temperature for three principal directions

in the composite: warp, fill, and normal.

The present study focused on the epoxy-resin matrix materials used

in these two composites. Young's modulus was determined semicontinuously

between ambient and liquid-nitrogen temperatures, and internal friction

was determined at the two end-point temperatures.

Physical properties of a composite's constituents are essential to

understanding the composite's behavior based on micromechanical models.

Indeed, prediction of macroscopic properties from constituent properties

is one of the few really fundamental problems in the theory of composite

materials. And it remains the only such problem on which significant

progress has been made.^

Experimental

a. Materials

Samples of each epoxy resin were obtained from a commercial source

that also produces the composite materials. These samples were processed

37n
in terms of temperature, pressure, and time as though they were being

fabricated into a composite. Both materials were translucent, the

G-10 form being light-amber colored and the G-11 form dark-amber. Mass

3 densities determined hydrostatically were 1.183 and 1.234 g/cm ,

respectively.

b. Methods

Methods used to measure Young's modulus and internal friction were

described previously by Ledbetter.^ Briefly, cylindrical specimens

5 mm in diameter and about 2 cm long formed one component of a a threecomponent (Marx) oscillator. The other two components were identical

x-cut quartz piezoelectric crystals with fundamental resonance frequencies near 45 kHz. Young's modulus, E, is determined from

E = 4pf^Jl^ (1)

where p denotes mass density, f frequency, i length, and subscript-s

specimen. The specimen's resonance frequency derives from

f = f + (f - f m^/m^ (2)

S 0 0 q' q' s

where m denotes mass, subscript-o total oscillator, and subscript-q

quartz

Internal friction, Q~\is determined by

Q"' = (f2 -

371
the resonance-peak width at 1//^ of rnaximum in a strain-versus-frequency

curve. Similar to Eq. (2),

^ (5;' - Qqh (4)

Resul ts

Room- temperature and liquid-nitrogen-temperature E and Q

^ values

are given in Table 1. Young ' s-modul i-versus-temperature results are

given in Table 2 and shown in Figs. 1 and 2.

Table 1. Young's modulus and internal friction at room temperature

and liquid-nitrogen temperature for two epoxy resins.

Young 's liodulus

(GPa)

Internal Friction

295 K 76 K 295 K 76 K

G-IOCR 3.80 8.37 0.0396 0.0057

G-llCR 4.49 6.55 0.0164 0.0032

372
Table 2. Temperature dependence

in units of 10“ N/m^.

of Young's moduli of G-IOCR and G-llCR

(Preliminary results.)

epoxy resins,

T(K) G-IOCR G-llCR

300 .03640 .04319

290 .03972 .04454

280 .04300 .04588

270 .04623 .04720

260 .0'4940 .04850

250 .05251 .04977

240 .05555 .05102

230 .05852 .05224

220 .06140 .05343

210 .06419 .05458

200 .06688 .05570

190 .06945 .05677

180 .07191 .05780

170 .07423 .05878

160 .07639 .05970

150 .07839 .06055

140 .08022 .06134

130 .08184 .06205

120 .08325 .06267

no .08444 .06320

100 .08539 .06364

90 .08611 .06397

80 .08661 .06421

76 .08676 .06429

373
2

. 00 h

V

-I

140

LTJ cn

CNI

o

>- UJ cu

Ln

CD

CNI

374

100

200

300

^°°0

200

300

TEMPERATURE

(K)

TEMPERATURE

(K)

Fig.

1.

Young's

modulus

versus

temperature

Fig.

2.

Young's

modulus

versus

temperature

for

G-IOCR

epoxy-resin

matrix

material.

for

G-llCR

epoxy-resin

matrix

material.
Di scussi on

As shown in Figs. 1 and 2, the Young's moduli of the epoxy resins

change regularly with temperature: linear at high temperature; negative

slope, decreasing with decreasing temperature; approaching zero slope at

zero temperature. Such regular behavior indicates the absence of phase

transitions or other abrupt changes in the internal structure of the

material. Such behavior can be described by various models that have a

3 theoretical basis. Varshni proposed that elastic-stiffness constants,

C, change with temperature according to

C(T) = C(o) - s/[exp(t/T - 1)1 (5)

where C(o), s, and t are the three adjustable parameters, and T is

absolute temperature. Physically, C(o) is the zero-temperature elastic

constant; t corresponds to a characteristic vibration temperature such

as the Einstein temperature; and -s/t is the high-temperature limit of

the thermoelastic coefficient dC/dT, the slope of the C(T) curve.

Parameters determined by fitting the observations to Eq. (5) are

given in Table 3. Consistent with a higher dC/dT, the G-IOCR epoxyresin shows a higher value of s, despite having a similar t. This

irregularity suggests that the two materials are structurally dissimilar. Some study may be appropriate to determine the physical

meaning of s, which determines the relative temperature behavior of

these two materials, a behavior usually determined by t.

The quantitatively different temperature behavior of the two

materials is dramatic. Expressed as -(1/E) (dE/dT) at room temperature,

375
Table 3. Parameters in equation (5).

C(o) s t (IQii N/m2) (loll N/m2) (K)

G-IOCR 0.0872 0.1926 470

G-llCR 0.0645 0.0726 445

the values are 0.0087 for the G-IOCR resin and 0.0031 for the G-llCR

resin, a difference of a factor of three' In terms of a linear-chain

model of solids and the temperature dependence of the elastic constant

r

• 5 Cii IS

dCii/dT = -g2k/2f2a (6)

where k denotes Boltzmann's constant, a the lattice parameter, and f and

g the second-order and third-order force constants. Since the similar

values of E show that f is about the same for the two materials, g must

be much higher for the G-IOCR case. This higher anharmonicity implies

that properties such as thermal expansivity and specific heat will be

higher also.

The dissimil arity of the two materials manifests itself also in the

internal friction, where the G-IOCR value is approximately double that

4

of the G-llCR case. For a large number of composite materials, Ledbetter

established an approximately hyperbolic inverse relationship between E

and Q"^ . The result results suggest that this relationship may be more

valid at ambient temperatures than at lower temperatures.

376
References

1. H. M. Ledbetter, Dynamic elastic modulus and internal friction in

G-lOCR and G-llCR fi bergl ass-cl oth-epoxy composites, in Material

Studies for Magnetic Fusion Energy Applications at Low Temperature- 11

,

NBSIR 79-1 609 (1 979), pp. 437-451.

2. Z. Hashin, Theory of composite materials, in Mechanics of Composite

Material

(Pergamon, New York, 1970), pp. 201-242.

3. Y. P. Varshni, Temperature dependence of the elastic constants,

Phys. Rev. 3952-3958 (1 970).

4. H. M. Ledbetter, Dynamic modulus and damping in fiber-reinforced

composites, to be published.

5. G. Leibfried, Gi ttertheorie der mechanischen and thermischen

Ei genschaften der Kristalle, in Handbuch der Physik , Volume 7, Part 1

(Springer-Verl ag , Berlin, 1955), pp. 104-324.

377
I

^ **!*»- 'twAi ••u^ "'

.v ^WT' A.«*'|Ji)^Jjv|i^

.'.•.>>» “ 'Vi • j.

.t’-i|S,«.'

, vi>e,
DIRECTIONAL VARIATION OF YOUNG'S MODULUS AND

INTERNAL FRICTION IN TWO FIBERGLASS-CLOTH-EPOXY COMPOSITES

National Bureau of Standards

379

DIRECTIONAL VARIATION OF YOUNG'S MODULUS AND INTERNAL FRICTION IN TWO

FIBERGLASS-CLOTH-EPOXY COMPOSITES

H. M. Ledbetter

Fracture and Deformation Division, NBS, Boulder, Colorado 80303

Abstract

Young's modulus was measured dynamically (near 50 kHz) for various

directions in the warp-fill plane and in the warp-normal plane of

laminated fiberglass-cloth-epoxy composites. Results are interpreted

using an anisotropic elasticity model. Directional dependence of

internal friction was determined on the same specimens. Directions of

high modulus are directions of low internal friction, and vice versa.

Key Words: Composites; elastic constants; glass-epoxy; internal friction;

Young ' s modul us

381
Introduction

Stimulated largely by aerospace applications, interest in composite

materials has skyrocketed during recent decades. The longstanding

interest in lightweight, high-strength, low-cost materials has also

intensified. For many structural applications, fiberglass-cloth-epoxy

composites provide these desired properties. These considerations

combined with el ectrical -property requirements led the National Electrical Manufacturers Association (NEMA) to develop specifications for

two materials: G-10 and G-11, with the corresponding cryogenic grades

G-IOCR and G-llCR.

A principal character! Stic of composites is the directional

variation of most of their properties, mechanical and physical. Indeed,

the primary theoretical problem for composites is to predict their

direction-dependent properties from the (usually isotropic) properties

of their constituents.

The macroscopic-microscopic elastic constants of composites have

been studied extensively for fiber-reinforced materials. This subject

1 2 has been reviewed by Bert from an experimental viewpoint and by Sendeckyj

from a theoretical viewpoint.

The present study determined the directional variation of both

Young's modulus and internal friction in two fiberglass-cloth-epoxy

composites: NEMA grades G-IOCR and G-llCR.

Experimental

a. Materials

Table 1 summarizes the properties of the materials, which were

received from a commercial source in the form of plates 2.5 to 5.0 cm

thick . Specimen lengths varied from 2 to 5 cm depending on material.

382
fiber orientation, and resonance frequency. Three or more specimens

were studied for each material and orientation.

b. Young ' s-modu1 us measurements

3 Experimental methods were described previously in detail elsewhere .

Briefly, the specimen formed one component of a three-component (Marx)

oscillator. The other two components were identical x-cut quartz

piezoelectric crystals with fundamental resonance frequencies near

45 kHz. Young's modulus, E, was determined from

E=4pf2i2

where p denotes mass density, f frequency, £ length, and subscript-s

/

specimen. The specimen's resonance frequency was derived from

.2 .2 ^ /.2 . 2 ^

, f =f +(f -f)m/m s 0^0 q^ q s ( 2 )

where m denotes mass, subscript-o total oscillator, and subscript-q

quartz.

c. Internal -friction measurement

Internal friction, Q~\ was determined by

= (f2 - q)/fo (3)

the resonance-peak width at 1//2 of maximum in a strain-versus-frequency

curve. Similar to Eq. (2),

= Q'^

+ (Q"^ - Q"^) m /m /.x

^s ^0 '^0 ^q q s (4)

Results

Figures 1-4 and Table 2 show the principal results of the study:

the directional variation of Young's modulus and internal friction for

383
two fiberglass-cloth-epoxy composites. Because of the similarity of the

two composites, complete sets of measurements were not made on both

materials. G-IOCR was studied in the warp-fill plane and G-llCR was

studied in the warp-normal plane. Samples were not available for

studying the fill-normal plane; but because the fiber density does not

differ markedly in the warp and fill directions, the warp-normal plane

should not differ markedly from the fill-normal plane.

Discussion

It is attractive to interpret the strong dependence of Young's

modulus on direction (Figs. 1 and 2) by invoking well-known anisotropic4 elasticity relationships . We assume tetragonal elastic symmetry, that

is, the equivalence of the warp and fill directions. Of course the

composites actually have orthorhombic elastic symmetry; but this means

nine independent elastic constants and a more complicated analysis with

less useful results. Results in Table 2 and previously reported results'

for the temperature dependence of Young's modulus show that the fillwarp equivalence is not too bad an assumption for present purposes.

Tetragonal -symmetry materials have six independent elastic constants. For the elastic compliances, the S— 's, these are usually taken

to be. ^11 ~ ^22^ ^33^ ^12^ ^13

~ ^23* ^44

~ ^55^ ^nd ^ 00

* Young s

modulus along the x-j axis is S^| , along it is > and along x^ it

is Herein, x-j , x^, and x^ correspond to the warp, fill, and normal

directions, respecti vely

For an arbitrary direction in a material. Young's modulus is

^11 ^1111 = “li“l j“lk“u^ijk£ (5)

334
where primes denote a coordinate system rotated away from x^, x^,

the four indices on S are the full -tensor form necessary for tensor

transformations; and the 's are the direction cosines between the

primed and unprimed coordinate systems. For tetragonal symmetry,

Eq.(5) becomes

S.J-J = (a^ + a^) S-j-| + + 2a^a2S-|2 + 2{a^a^ + a2a3)S-|2

, 2 2 2 2,^ ^ 2 2c + [a20t-2 ( 6 )

For directions in the x-|-X

2 plane, Eq. (5) becomes

S'li ( a-| + a

2

) S

^^ + a ^«2 ( 25^2 (7)

For directions in the x-j-X

3

plane, Eq. (5) becomes

S^l + a3S23 + a-^a^ (2S^2 ^ 44' ( 8 )

which is by symmetry the same as for the X

2 -X3 plane when (X 2 Is substituted for a-j

From Eqs. (7) and ( 8 ) and the present experimental results, leastsquares values of 25.^2 Sgg and 25-^2 + can be computed. The results

are 16.05 and 16.37, respectively. From the relationship

$12 -v-|2S-|^ (9)

assuming v-j 2

= 0.165 , S-j 2 is -0.5593. Thus, Sgg is 17.17 and

^66 ~ ^66 ~ 0-058. Cgg represents the resistance on the X

2 or x-j plane

to a shear stress directed along either x-j or X

2

* Consistent with the

results in Fig. 1, this result indicates the lower rigidity of the

composite to any deformation not directed uniaxial ly along either

x.| or X

2

*

385
Similarly, from the relationship

S 13 ''13^11 ( 10 )

an(i the assumption that v-j^ = 0.333, a typical value for an epoxy resin,

which constitutes the matrix, which largely determines the v-j^ value, it

results that S-j^ = -0.105 and that = 16.58. Thus = 0.060

is the resistance on the x^ or x^ plane to a shear stress along either

x^ or x^. As for the case of a C^g-type shear, this deformation strongly

reflects matrix properties and has a low value.

Thus, assuming tetragonal symmetry, from measurements of the Young 'smodulus directional dependence and two assumptions concerning Poisson's

ratio, we have deduced the complete set of six elastic constants of a

fiberglass-cloth-epoxy composite. These are listed in Table 3. At the

same time, the seemingly irregular directional dependence of Young's

modulus is explained as resulting from contributions from elastic constants

such as and Interestingly, for this material, the elastic

constants S-j 2 and S-j^ affect Young's modulus only negligibly.

Internal friction shows an inverse relationship to Young's modulus,

as shown by comparing Figs. 3 and 4 with Figs. 1 and 2. This relationship

holds for a wide variety of composites, as "shown by Ledbetter”. While

one intuitively associates higher damping with softer elastic stiffness,

the author knows no model that shows this explicitly. Similarly, the

author knows no model for interpreting anisotropic internal friction.

386
Concl usions

The following principal results and conclusions emerge from the

present study:

1. Young's modulus, E(a^. )> varies strongly with direction,

both in the warp-fill plane and in the warp-normal plane.

2. This directional variation can be interpreted using an

anisotropic elasticity model.

3. Alternatively, by making two assumptions concerning

Poisson's ratio, the complete set of elastic compliances, the S^.^.'s can

be determined from the E(a^. ) results.

4. Internal friction varies inversely with E(a^). This

result, while not unexpected, cannot be interpreted or explained by any

simple model known to the author.

Acknowl edgment

Both R. E. Schramm and M. B. Kasen of NBS provided information and

materials required for this study.

387
References

1. C. W. Bert, Experimental determination of composites, in Composite

Materials, Volume 8 (Academic, New York, 1 975).

2. G. P. Sendeckyj, Elastic behavior of composites, in Composite

Materials, Volume 2 (Academic, New York, 1974), pp. 45-83.

3. H. M. Ledbetter, Dynamic elastic modulus and internal friction in

fibrous composites, in Nonmetal lie Materials and Composites at Low

Temperatures (Plenum, New York, 1979), pp. 267-281.

4. J. F. Nye, Physical Properties of Crystals (Oxford Clarendon,

London, 1 960), p. 131-149.

5. R. E. Schramm and M. B. Kasen, Fracture and Deformation Division,

NBS, Boulder, Colorado, personal communication (1979).

6. H. M. Ledbetter, Dynamic elastic modulus and damping in fiberreinforced composites, to be published.

388
Table 1. Properties of studied materials.

Mass Density Fiber Volume

Material Fi ber Matrix (g/cm3) Fraction

G-IOCR 7628 E-glass cloth bisphenol epoxy 1.90 0.64-0.68

G-llCR 7628 E-glass cloth bisphenol A, aromatic amine 1.95 0.64-0.68

Table 2. Dynamic Young's moduli, E, and

of G-IOCR and G-llCR.

internal friction, Q~ 1

5

Temp Di recti on G-IOCR 1 A G-llCR 1 A (K) E(GPa) Q'^IO’"^) E(GPa) Q'^10"^)

295 warp 29.4+1 .0 68.3±6.6 31 .4+0.4 147.0+34.9

fill 26.3+1.0 114.0+7.8 27.7 161.3

normal 14.0 228.6 15.6+0.6 406.5

76 warp 34.4 14.6 37.6 33.4

fill 32.9 17.3 33.6 -- normal 21.8 19.6 22.8 --

76/295 warp 1.17 0.21 1.20 0.23

dimensionless

ratios fill 1.25 0.15 1.21 -- normal 1.56 0.09 1.46 —

389
Table 3. Elastic compliances, the

elastic stiffnesses, the C..

for G-IOCR case.

's, and elastic stiffnesses, and

's, reduced as described in text

s.C. . ij

ij (10"^^m^/N) (lO^^N/m^)

11,22 3.67 0.279

33 6.41 0.156

44,55 16.58 0.060

66 17.77 0.058

12 -0.56 0.043

13,23 -0.11 0.005

390
Figure Captions

1. Directional variation of Young's modulus of G-IOCR composite in

warp-fill plane, in units of GPa.

2. Directional variation of Young's modulus of G-llCR composite in

warp-normal plane, in units of GPa.

3. Directional variation of internal friction, Q”\ of G-IOCR com- posite in warp-fill plane,

4. Directional variation of internal friction, Q~\ of G-llCR com- posite in warp-normal plane. Two points at 15° and 30° represent

specimens from different samples of composite. Dashed line indicates

that the exact directional variation remains uncertain in this

region.

391
Warp

Warp

CDO

392

Directional

variation

of

Young's

modulus

2.

Directional

variation

of

Young's

modulus

of

G-lOCR

composite

in

warp-fill

plane,

6-llCR

composite

in

warp-normal

plane,

in

units

of

GPa,

in

units

of

GPa.
iA

0) \ I

raEo

cycoo

•I—JfOc

S0) 4-> c

fCua i § " ^ a;

M- 5 co

03

i-m>

fOcou

O) . 00

O)

C

<“ ^ ^ I— ro 03

fO 00 ^ E^ . c S ^ ^ o S &|

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393

THERMAL CONDUCTIVITY OF G-IOCR AND G-llCR INDUSTRIAL

LAMINATES AT LOW TEMPERATURES (PART II)

National Bureau of Standards

395

THERMAL CONDUCTIVITY OF G-IOCR AND G-11CR INDUSTRIAL

LAMINATES AT LOW TEMPERATURES (PART II)

J. G. Must

ABSTRACT

Experimental thermal conductivity results are reported for G-IOCR composite

specimens from two suppliers and at strains up to 90% of ultimate at temperatures from 2 to 300 K. The results indicate considerable differences between

the two manufacturers but essentially no effect due to strain. Data are also

reported on the resins used in G-IOCR and G-llCR. These results are difficult

to understand and further measurements are scheduled to define their properties

more clearly.

INTRODUCTION

NBS is performing extensive 1 ow- temperature electrical, thermal, and mechanical properties characterization of cryogenic grades of G-10 and G-11 fiberglass-epoxy composites. These special laminate grades, referred to as G-IOCR

and G-llCR, meet NEMA Grades G-10 and G-11, MIL-P-18177C, and federal specification LP-509.

In a previous report (Part I) [1], thermal conductivity data were presented

for G-IOCR and G-llCR at temperatures from 2 to 300 K for heat flow in both the

normal and warp directions. These data indicated that the two laminates, produced by a single manufacturer (manufacturer A), had the same conductivity to

within about 10%. Also, it was found that the warp direction conductivity was

about 20% higher than that in the normal direction.

The purpose of the present work is to investigate the difference in thermal

conductivity on specimens obtained from another manufacturer (manufacturer B),

to investigate the effect of large strain, and to determine the thermal conductivity of the pure epoxy resin used in these laminates. The latter results will

aid the analysis and interpretation of the composite data.

397
MATERIAL AND SPECIMEN CHARACTERIZATION

Both G-IOCR and G-llCR are woven glass-epoxy composites. G-IOCR is a heatactivated amine-catalyzed bisphenol A, sol id- type, epoxy resin laminate reinforced with continuous filament E-glass fiber fabric, silane finished. G-llCR

is an aromatic amine bisphenol A liquid- type epoxy resin laminate, reinforced

with the same fabric. The fabric thread count is 17 1 per cm (43 3 per in)

for warp and 12 ^ 1 per cm (32 2 per in) for fill. The threads are composed

of 75 filaments each of 0.00046-cm diameter (0.00018-in). The direction referred to as normal in these measurements is perpendicular to the cloth plane.

The density of E glass is 2.4 g/cm^. The resin used in G-IOCR has an

uncured density of 1.10 0.01 g/cm^, and the G-llCR resin has an uncured density of 1.04 +_ 0.01 g/cm^. The G-IOCR composite contains 33 to 35 weight percent resin, and the G-llCR contains 27 to 31 weight percent resin. The typical

density of G-llCR is given by the manufacturer as 1.97 g/cm^.

The thermal conductivity specimens of G-IOCR and G-llCR lamintes are formed

by cutting approximately 2-cm squares from a laminate sheet 4-mm thick and

cementing these squares together to form a cube. The squares were cemented

using a strain-gage adhesive under pressure at 150°C to obtain a thin glue line.

The cured blocks were then machined to the dimensions shown in Table 1. The

specimen masses and calculated densities are also given.

The resin specimens were cut from larger blocks supplied by the manufacturer. The dimensions, masses, and calculated densities are given in Table 1.
TABLE 1 DIMENSIONS, VOLUMES, WEIGHTS, AND DENSITIES

OF THE THERMAL CONDUCTIVITY SPECIMENS

SPECIMEN X^^1

DIMENSIONS

(cm)’

X^ X^^3 VOLUME

( cm)^

WEIGHT

(g)

DENSITY

(g/cm)

G-10CR/WARP(78) 1 .910 1 .890 1.919 6.93 13.188 1 .904

G-11CR/WARP(78) 1 .915 1.913 1 .910 7.00 13.684 1 .956

G-10CR(6)

WARP (90 PCT)

1.913 1 .911 1.916 7.00 12.767 1 .824

G-10CR(1

)

WARP (As Rec.)

1.916 1.819 1 .824 6.36 11 .608 1 .825

G-IO(GW) NORM 1.882 1.908 1 .873 6.73 11 .717 1 .741

G-IOCR RESIN 1 .915 1.913 1 .781 6.52 7.698 1 .181

G-llCR RESIN 1 .915 1.915 1.915 7.02 8.620 1 .228

X] is normal to the fabric for the composite specimens

X

2 is the warp direction for the composite specimens

X

3 is the fill direction for the composite specimens

309
Using the measured density of each specimen, the given density of E glass,

and the weight fractions of the resins in the composite specimens, the resin

densities can be calculated with the equations below:

Pc =Pc WFr '^''g

^c = density of composite

^r = density of resin

^g = density of glass

WFp = weight fraction of resin

WFg = weight fraction of glass

This resulted in a resin density of 1.25 0.01 g/cm^ for both resins. Note

that the measured densities of the two resin specimens are 1.181 g/cm^ for

G-IOCR and 1.228 g/cm^ for G-11 CR resin. By using the equation

Vc = VF^ + Vg VFg

where V is specific volume (V = 1/p) and VF is volume fraction, the volume fiber

content was calculated to be 53% for G-IOCR and 58% for G-llCR.

EXPERIMENTAL PROCEDURE AND DATA ANALYSIS

The thermal conductivity of the test specimens is determined using a

thermal conductivity integral apparatus. This apparatus is described in detail

by Hust and Arvidson [2]. The estimated uncertainty of the results is less than

10%, unless otherwise noted, as verified by measurements on NBS standard

reference material, SRM 735.

400
RESULTS AND DISCUSSION

The specimen designated as G-10CR(6) WARP (90 PCT) was prepared by first

cutting a 2-cm strip from a sheet of material that was room-temperature strained

in the warp direction to 90% of ultimate. This strained strip was then cut into

pieces and glued together to form a block, as previously described. Seventeen

data points (thermal conductivity -integral s) were obtained at temperatures from

4 to 304 K. Heat flow was in the warp direction. These logarithmically spaced

data were represented by the thermal conductivity integral method (see Hust and

Arvidson [2]) to yield the parameters of the following equation:

n K(T)= ^ A [ln(T + 1)]^ (Eq. 1)

i = l

The rms relative deviation of the experimental values of thermal conductivity

integral from the least squares fit is 1.0% with a maximum deviation of 3.6%.

Values of K(T) are illustrated in Figure 1.

It is noted that these values of K(T) differed significantly from the data

reported in 1978 for an unstrained specimen of G-IOCR/WARP. The data for

G-IOCR/WARP (78) and G-1 1 CR/WARP( 78) are repeated in Figure 1 for comparison.

To determine if this difference was caused by material variability (the two sets

of specimens were prepared from sheets of materials supplied by different manufacturers) or by the strain, another specimen was prepared from the same unstrained sheet as the above strained specimen. This unstrained specimen is

identified as G-lOCR(l) WARP (as rec.). It was measured only at 281 and 88 K.

These data are plotted on Figure 1 and reveal that the effect of strain is

401

undetectabl e.
M ‘AllAllOnaNOO TVlAld3Hl

.'oz

TEMPERATURE,
Because of the large material variability suggested by the resul.ts of the

1978 and 1979 measurements, as well as by unpublished data on similar materials,

further measurements on stock G-10 were undertaken. A specimen was fabricated

from off-the-shelf 0.95-cm thick G-10 sheet. This specimen, referred to as

G-IO(GW) NORM, was fabricated as described above and measured at 287 K with heat

flow normal to the fabric. The thermal conductivity is lower by 19% than the

value obtained in 1978 on G-1 OCR(NORM) . It appears from this limited data that

the CR grade of G-10 has a higher conductivity than the standard G-10 material.

To facilitate future analysis of these and other composite data, measurements were performed on specimens fabricated from the pure resins used in G-1 OCR

and G-llCR. These specimens were prepared from larger cured pieces supplied by

manufacturer A. Twenty-one data points between 2 and 304 K were obtained on

G-1 OCR Resin and sixteen data points were obtained from 4 to 304 K on G-1 ICR

Resin. The results, identical within the imprecision of the data (+^3%) for both

resins, are illustrated in Figure 1. These values are more uncertain than the

previous results because the specimen thicknesses were too large for optimum

accuracy. The presence of heat transfer by radiation through the specimen is

also potentially a problem in defining unique values of thermal conductivity.

That is, there may be a measurable thickness effect. Further work has been

scheduled to define the resin values more clearly. The present estimated uncertainty for the smoothed resin data is + 15%.

403
REFERENCES

1. M. B. Kasen, J. G. Hust, and H. M. Ledbetter, "Nonmetal 1 ics for Magnet

Systems," NBSIR 79-1609, pp. 405-451 (1979).

2. J. G. Hust and J. Arvidson, "Thermal Conductivity of Glass Fiber/Epoxy

Composites Support Bands for Cryogenic Dewars," Unpublished report

275.03-78-2.

404
THERMAL EXPANSION OF CRYOGENIC-GRADE GLASS-EPOXY LAMINATES

National Bureau of Standards

405

THERMAL EXPANSION OF CRYOGENIC-GRADE GLASS-EPOXY LAMINATES

M. A. Ranney and A. F. Clark

ABSTRACT

Thermal expansion measurements for special grades of NEMA-type G-10 and

G-11 glass-epoxy laminates were carried out in a quartz-tube dilatometer over a

temperature range 4 to 300 K. It was found that expansions in the warp direction were comparable to those of the samples' reinforcement fabric, whereas

expansions in the normal direction were characteri stic of the samples' epoxy

resins

INTRODUCTION

Special grades of NEMA-type G-10 and G-11 glass-epoxy laminates (G-IOCR and

G-llCR) were measured for thermal expansion to investigate further their use in

cryogenic properties.

These laminates are commercial products meeting current NEMA G-10 and G-11,

MIL-P-18177C, and federal LP-509 specifications. G-IOCR is a heat-activated,

amine-catalyzed bisphenol A sol id- type epoxy resin laminate reinforced with

continuous-filament E-glass fabric, silane finished. G-llCR is an aromatic-amine

hardened bisphenol A liquid- type epoxy resin laminate reinforced in the same

manner.

Both laminates are woven glass cloth composites. The warp fibers are those

which emerge from the loom, and fill fibers are those woven across them at right

angles. Each composite has a yarn count of 17 1 per cm (43 3 per in) for

warp and 12.6 + 0.8 per cm (32 _+ 2 per in) for fill. The composites' normal

direction is perpendicul ar to the cloth ply plane.

Our measurements for G-IOCR and G-llCR are in the warp and normal directions.

407
SAMPLE PREPARATION

For measurements in the normal direction, samples were prepared using eight

pieces of each laminate, which were bonded end-to-end by a commercially available epoxy resin adhesive. This adhesive has special low expansion properties

and was used to minimize our experimental error. The G-IOCR sample dimensions

were 6.5 rnn x 6.5 mm x 19.92 mm, and the G-llCR sample dimensions were 6.5 nm x

6.5mm x 19.76 mm.

The samples for warp direction measurement were both directly machined to

the dimensions 4.2 mm x 4.2 mm x 20.32 mm.

APPARATUS

Measurements were made in a quartz tube dilatometer [1], schematically

shown in Figure 1. The differential contraction is transmitted to a room temperature mechanical dial gauge (sensitivity 0.5 ^m) by concentric quartz tubes

(labeled quartz support tube and quartz rider in Figure 1). All samples were

kept vertical by a copper heat shield, which also inhibited thermal gradients

from interfering with the measurement accuracy.

Samples were each cooled at a rate of 2.5 K min~^ from 300 to 4 K, and

the temperature was determined by a chromel -constantan thermocouple attached

directly to the samples at midpoint. The helium gas temperature and the sample

temperature were assumed to be the same.

The apparatus was calibrated by measuring the thermal expansion of a polycrystalline OFHC copper bar. The difference between the average measured data

and the standard reference material (SRM) data for OFHC copper [2] was used as

an absolute correction factor. Measurement accuracy was estimated at + 0.005%.

408
RESULTS

Figure 2 shows the relationship of the four samples tested (G-IOCR warp and

normal, G-llCR warp and normal), graphically depicted in terms of thermal expansion.* A computer curve fitting program was developed to fit the experimental

data, [L(293) - L(T)]/L(293) , and to approximate expansion coefficients,

[1 /L(293)] (dL/dT) K"l, from the data. The form of the function was

The data obtained indicate that thermal expansion in the warp direction is

comparable to that of the composites' E-glass reinforcement fabric, while expansion in the normal direction is characteristic of the bisphenol A epoxy resin.

Aside from the present study, only one other reference of G-IOCR and G-llCR

expansion data was located. These data on the same material , independently provided by the Los Alamos Scientific Laboratory, R. I. Schermer [31, gave the

total contraction from 295 K to 75 K using a fixed-point apparatus. The Los

Alamos data is found to be consistently ~10% greater in total contraction than

the data in the present study, yet the difference is within the accuracy limits

of the two sets of measurements.

*The SRM 736 Copper Standard is included as a scaling reference.

1. A. F. Clark, Cryogenics 8, 282 (1968).

2. T. A. Hahn, J . Appl . Phys . 41 , 5096 (1970).

3. M. B. Kasen, J. G. Must, and H. M. Ledbetter, "Nonmetal 1 ics for Magnet

Systems," in Materials Study for Magnetic Fusion Energy Applications at Low

Temperatures-I I

,

NBSIR 79-1609 (1979), p. 405.

L(293) - L(T>

^ g ^ L(293)

( 1 )

Table 1 summarizes the data for the four samples.

CONCLUSIONS

REFERENCES

409
Bell Jar

Figure 1. Schematic representation of quartz tube dil atometer.

410
THERMAL

EXPANSION

10

Figure 2. Thermal expansion of G-IOCR and G-11CR in both

warp and normal directions.

411
Table

1.

THERMAL

EXPANSION

DATA

FOR

G-10CR

and

G-llCR

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^;-12
Technology Transfer

STRUCTURES, INSULATORS AND CONDUCTORS

FOR LARGE SUPERCONDUCTING MAGNETS

National Bureau of Standards

U. S. Department of Energy

STRUCTURES, INSULATORS AND CONDUCTORS FOR LARGE SUPERCONDUCTING MAGNETS*

F. R. FICKETT and R. P. REED

National Bureau of Standards, Boulder, CO 80302

i

I and

! E. N. C. DALDER

Department of Energy, Washington, D. C. 20545

1

1

I As the size of superconducting magnet systems increases, the structural aspects become of

I of Increasing Importance. Because the stresses increase rapidly with size and field, magnets

i of the type proposed for new devices will be largely structure. We are engaged in a program

to acquire low temperature mechanical and physical property data on a wide range of mater-

, lals. The present status of this data base is reviewed and results from our test program are

! described. We have placed emphasis on the austenitic stainless steels as the primary struc- tural material, although data have also been acquired on a few aluminum alloys. Particular

I stress has been placed on the fracture characteristics of both the base metal and weldments

j' where appropriate. Achieving a good match between weld and base metal at 4 K is proving to be a difficult task. Several industrial fiberglass-epoxy laminates of fixed composition have

. been produced as a result of this program. These materials, GIO CR, Gil CR and Gil CR(BF)

have been tested at 4 K and the results are presented. Recent measurements on the effect of stress and fatigue on matrix materials for superconductor stabilization are also briefly re-

' viewed. A new handbook of data on the properties of structural materials for superconducting

' magnet systems is described.

1. INTRODUCTION

It now seems nearly certain that economical

fusion reactors relying on magnetic confinement

will require superconducting magnets, at least

for the toroidal field colls and probably for

other coll systems as well. Such magnets operate

well below 20 K and most often very near to 4.2

K, the temperature of liquid helium at one at- mosphere. There is essentially no chance that a superconductor capable of operation at significantly higher temperatures will be developed and

brought to the commercial stage in time to do us any good. The question of which of the existing

superconductors will be adequate for fusion en- ergy applications remains to be answered. NbTl

is certainly the most preferable since it is a ductile alloy and, thus, easily fabricated and

quite strain tolerant, but such magnets are re- stricted to maximum fields of about 8 tesla.

The intermetalllcs Nb 3 Sn and V 3 Ga will handle much higher fields, but are expensive to produce and very sensitive to strain.

Regardless of the superconducting material

finally chosen, the magnets will be large by

present-day standards. Most of the magnet will

be structure necessary to confine the large

forces generated by the field. At the same

time, problems of access to the plasma region

severely limit the total amount of structure

which can be tolerated in a working device.

Alloys with relatively high strength are a clear

choice here. Furthermore, the structure must be

at cryogenic temperatures near 4 K, a region

*Sponsored by the Office of Fusion Energy,

Department of Energy, Washington, D.C. 20545

where little information exists on material

properties. This lack of data is especially

serious in the case of welded joints between

large structural members.

Working superconducting magnets also require

large amounts of nonmetallics as electrical

insulation and as thermal standoffs. Where

rapidly changing magnetic fields are encountered, it is frequently desirable to make the en- tire structure from nonmetallics in order to avoid heat input due to eddy current generation.

Furthermore, the multilayer "super insulae ion" now favored for dewar construction is a metallized polymer film also subject to eddy current heating. The properties of many of these materials are not well characterized for low temperature service. Unlike alloys, they are not often

of standard composition and, even when they are, material variability is high. The neutron and

gamma radiation which will be found in working

reactors Introduces a further complication in the use of nonmetallics. It is not at all certain

what the useful lifetimes of materials will be under the combined Influence of low temperature

and radiation.

The final major magnet component is the con- ductor composite. The superconducting material

is just one part of it. The composite also con- tains a large amount of pure metal, usually

copper, to protect the superconductor in the event of an excursion out of the superconducting

state bv carrying the current for a time long

Presented at the First Topical Meeting on Fusion Reactor Materials, Jan. 1979,

to be published in Journal of Nuclear flaterials 1979.
enough to allow the superconducting state to be reestablished. The resistivity of this stabilizing material at the operating temperature is a critical parameter and one which may be strongly

Influenced by stress, fatigue and radiation.

Most of the problems outlined here were recognized a long time ago, but only recently has

there been a concerted effort toward their solution. In this paper^ we attempt to outline the

progress of that effort and to summarize the

present status. For many readers, this world of

low temperature material behavior is somewhat

new, and thus, our presentation will be, in part,

tutorial. The details will be found in the

references and the actual data in the various

handbooks and monographs described here.

A comment should be made as to how the low

temperature materials problems encountered in fusion energy applications differ from those

which occurred in earlier applications, such as the space program. Here very long term reliability will be required. The use of large sections

of material at low temperatures is new. The need

to weld these large pieces occurred only rarely

in the past. The projected amount of material is very large and, thus, necessitates production by

conventional mill techniques. On- site construction may be necessary. Complex NDE may be needed. Economics is a large factor in the materials

choices

2. STATUS OF THE LOW TEMPERATURE DATA BASE

The availability of a substantial data base

covering the mechanical and physical properties

of all materials used in magnets for fusion energy should be a matter of prime concern. It is important to stress that many materials behave

much differently at very low temperatures than

they do at room temperature. Extrapolation of room temperature properties data can lead to very

serious errors with often catastrophic conse- quences .

The data base as it exists today is barely

adequate in a few Instances and totally inadequate in most others. In no case does a sufficient data base exist to allow preparation of a code. Seldom is there enough to even permit a good assessment of the effect of material variability which, again, may be significantly different at low temperatures than at room temperature. It is fair to say that nearly no data

exist on the properties of alloy welds and of nonmetalllcs to 4 K. Data on fracture and

crack growth are just now being developed for most alloys of interest. Nevertheless, a few

general handbooks exist, as do monographs

treating specific properties and reports from

experimental test programs. Here we will describe the most common of these.

2.1 Data handbooks

The handbooks described here usually treat a

large number of materials and present many pro- perties. Most often data points are not shown,

and in many Instances, the curves displayed

are drawn with only limited data. Usually me- chanical properties data are taken at only a few discrete temperatures (4, 20, 77, 295 K)

and a complete curve is drawn. The most often

used handbooks which concentrate on low temper- ature data are (NTIS order numbers are in parentheses) :

a . A Compendium of the Properties of Ma t e- rlals at Low Temperatures, Part II, Pro- perties of Solids (1960). V. J. Johnson,

ed . (SS/PB 171619). Physical properties

(thermal) of metallic and nonmetalllc ma- terials.

b. Cryogenic Materials Data Handbook (1968).

F. R. Schwartzberg , ed . (AD 713 619 and AD 713 620). Mechanical and physical

properties of ferrous and nonferrous

alloys as well as plastics and nonmetallic composites.

c . Handbook on Materials for Superconducting

Machinery (1977). (ADA002-698 or Metals

and Ceramics Information Center

MCIC-HB-04). Mechanical and physical pro- perties of pure metals, ferrous and non- ferrous alloys, polymers and composites.

d . LNG Materials and Fluids (1977). D. B. Mann, ed . (Cryogenic Data Center, NBS).

Mechanical and physical properties of ferrous and aluminum alloys and thermal

insulation for tankage. Thermodynamic

properties of LNG components and nitrogen.

e. Materials for Superconducting Magnet

Systems (1979). (In preparation by NBS and Mechanical Properties Data Genter).

Mechanical and physical properties of

selected stainless steels and aluminum

alloys

In addition to these, many handbooks of data

at room temperature and above often carry some curves to the cryogenic temperature range.

2.2 Monographs

Several monographs are available which treat a specific low temperature property or present

data for a limited number of materials. Some of the more useful are:

a. R. J. Gorruccinl and J. J. Gnlewek,

"Specific Heats and Enthalpies of Technical Solids at Low Temperatures," NBS,

Monograph 21 (1960) (PB 172161).

b. R. J. Gorruccinl and J. J. Gniewek,

"Thermal Expansion of Technical Solids at Low Temperatures," NBS, Monograph 29 (1961) (PB 172178)
c. R. P. Reed and R. P. Mikesell, "Low Tem- perature Mechanical Properties of Copper

and Selected Copper Alloys," NBS, Monograph 101 (1967).

d. G. E. Childs, L. J. Ericks and R. L. Powell, "Thermal Conductivity of Solids at Low Temperatures. A Review of the Literature," NBS Monograph 131 (1973).

e. R. E. Schramm, A. F. Clark and R. P. Reed, "A Compilation of Mechanical, Thermal and

Electrical Properties of Selected Polymers," NBS Monograph 132 (1973).

2.3 Project reports

Two projects In recent years have attempted to acquire a wide range of data on structural materials at temperatures down to 4 K. The first was performed by NBS and Industrial contractors for

the Defense Advanced Research Projects Agency

(DARPA) and completed In 1977. It resulted In the handbook described above In Section 2.1(C).

The six semi-annual reports describe measurements on a number of alloys and a few nonraetalllcs. They Include all the details of apparatus

and procedure. The six volumes are available

under the title "Materials Research for Superconducting Machinery" (AD780596, ADA 004586,

ADA012365, ADA019230, ADA030170,. ADA036919).

The second project, also by NBS, Is for the Office of Fusion Energy of DOE and is currently in

its third year. It will be described in more detail shortly, but is essentially a program to develop a low temperature data base for selected

materials which have promise for use in superconducting magnets for fusion. The annual re- ports are titled "Materials Studies for Magnetic

Fusion Energy Applications at Low Temperatures."

Two are available from the authors of this pa- per .

2.4 Current awareness documents

The NBS Cryogenic Data Center publishes a biweekly listing of titles from the current

literature dealing with cryogenic research and

engineering. The title Is "Current Awareness

Service." The same organization publishes a quarterly, "Superconducting Devices and Materials." The Metal Properties Council and

Battelle publish a "Review of the Low-Temperature Properties of Metals" on a regular basis

which contains actual data extracted from the literature

3.0

MATERIAL BEHAVIOR AT LOW TEMPERATURES

As the temperature Is lowered, all mechanical

and physical properties of materials undergo

significant changes. Here we will look at the general behavior of these properties for some classes of materials. All of the discussions

will refer to Figure 1, In which we show schematic curves Illustrating the effect of tem- perature. The actual curves are typical of

austenitic stainless steels, but we will also

be discussing the behavior of pure metals and

nonmetallics . We want to stress that the

figures are only schematic. The behavior of any actual material should be checked In one of the handbooks described earlier.

3.1 Tensile properties (Figure 1 a,b,c,d)

The ultimate tensile strength of nearly all

materials increases as the temperature is lowered and the increase Is usually quite large.

It is not greatly affected by alloying additions such as nitrogen in the stainless

steels

The tensile yield strength also increases

but not as much. In the stainless steels the amount of Increase Is strongly dependent on nitrogen content with' high nitrogen giving a much larger increase. For the nonmetallic

composites the yield strength is about the same as the ultimate strength. Polymers show

a varied behavior and no generalization is possible

The Young's modulus also increases with decreasing temperature for most materials. The

increases are not particularly large but are significant for engineering design.

The elongation does not always decrease uniformly as one might expect, even for the polymers. Under the combined influence of temper- ature and stress, many materials undergo

transformations (e.g., martensite formation in stainless steels) which have a strong Influ- ence on elongation. Changes In elongation are seldom a problem in actual applications^ since

structures are never taken that close to fail- ure .

3.2 Fracture and fatigue (Figure 1 e,f,g)

These properties are very significant ones

for fusion applications, particularly where

pulsed field operations are encountered. They

are strongly affected by temperature and by

material composition. Little low teraperatiire

data exists for either pure metals or nonmetallics.

The fracture toughness may either increase or decrease with decreasing temperature. The desireable situation is to have an Increase, but the opposite is usually the case. There Is an Inverse correlation with the amount of increase

In yield strength, and thus, for example, nitrogen additions to stainless steel result In a fracture toughness decrease at low temperatures.

The low values found at 4 K may make many alloys

unacceptable for service.

The fatigue crack growth rate is another Im- portant property which can exhibit a variety of behavior depending on the material. This para419
meter is not well correlated with any other,

although It is obviously related to both yield

strength and toughess, but in a very complex

way. Again, the low temperature crack growth

properties may make material with promising properties at room temperature completely unacceptable .

Fatigue strength is a somewhat more well-behaved property. The high-cycle fatigue strength

of alloys and pure metals generally increases

with decreasing temperature. In stainless

steels a decrease is seen under conditions of

low-cycle fatigue (high stress levels). Non- metalllc composites show an Increased fatigue

strength as the temperature is lowered, but this

may be followed by a decrease below 77 K.

3.3. Thermal properties (Figure 1 h,i,j,)

In general, the physical properties are better

behaved than those already discussed although

wide variations are found, particularly in

"pure" metals, depending on the exact makeup of the metal.

The thermal expansion of essentially all

structural materials is negative, i.e., they con- tract as they cool. The rate of contraction de- creases toward lower temperatures leading to a strong decrease in the thermal expansion coefficient, a parameter not shown on the graph.

The thermal conductivity of most materials

decreases rapidly with lowering temperature.

Very pure metals and some dielectrics may show a large peak in the conductivity at low temperatures, but in all cases, the value at A K is well

below that at room temperature.

The specific heat of materials also decreases

very rapidly at low temperatures, usually dropping several orders of magnitude to 4 K. Unlike

the conductivity this property is quite insensitive to material composition. For example, the

specific heat behavior of nearly all austenitic

stainless steels is about the same. The very

low value of this parameter at 4 K can lead to serious problems of component heating, since a very small energy input can cause a large rise

in temperature.

(c) (d)

(e) (f)

T

(g) (h)

Thermal Expansion

Ll _ I T

Figure 1. The behavior of the mechanical and physical properties of materials as a function of temperature. The graphs are general schematics most representative of the austenitic

stainless steels.

^•20
3.4 Electrical and magnetic properties (Figure

1 k,l)

The electrical resistivity of metals always

decreases with decreasing temperature. The

total change observed In highly alloyed metals

Is quite small, usually less than a factor of

five. Pure metals, on the other hand, may show

decreases of several orders of magnitude. The

resistivity change tends to level out at the lowest temperatures. For organic nonmetalllcs

limited data Indicate that the volume resistivity Increases significantly and the dielectric

breakdown strength Increases slightly to 4 K.

The magnetic properties of most structural

materials are not much affected by low temperatures. The ferrous alloys are an exception.

In particular, the susceptibility of the aus- tenitic (nonmagnetic) stainless steels exhibits

a cusp presently thought to Indicate a spin

glass transition. The behavior of magnetization with field at low temperatures Is nonlinear and shows a tendency to saturation. Magnetic steels show pretty much the same behavior

as at room temperature, except that the saturation magnetization may be Increased a bit.

This discussion has been very brief of necessity. A more extensive treatment with many

specific examples Is provided by several earlier reviews [1,2,3]. In addition, a tutorial

book covering all aspects of materials behavior

at low temperatures Is now In preparation [4].

4. THE NBS/DOE MATERIALS RESEARCH PROGRAM

Two years ago the NBS Cryogenics Division

started a program for DOE (then ERDA) which had

as Its goals: a. the evaluation of proposed

fusion devices to determine the materials problems unique to the low temperature magnet systems; b. development of a research program to Insure that the necessary properties data would

be available; and, c. the rapid dissemination

of the data to the fusion community. As with

all programs, the available funding was limited

and, thus. It was necessary to assign rather

strict priorities to the tasks. Furthermore,

when the program started, there were no actual

superconducting fusion magnets or models under

construction^ so many of our decisions were based

on limited Information. The advent of the Large

Coll Program and the MFTF (Mirror Fusion Test Facility) project has resulted In some redirection of the research, but our Initial choices

have proven to be quite reasonable and the data

developed to date by NBS and Its contractors has

already been used extensively by the projects.

Here we will point out some of the more Inter- esting results and their Implications. The

total program Is described In the reports mentioned In Section 2.3. Our survey of fusion device designs and the research program proposals

which resulted are an appendix to the first re- port. This program specifically excluded re- search on the superconducting composite Itself.

Other NBS projects have measured properties of

some relatively small wires, however, and some results will be discussed In the next section.

The actual conductors to be used for fusion

magnets are very large and carry very high cur- rents (^ 10 kA at full field) and, thus, can only be tested In the few laboratories with

large magnet facilities.

The materials which we decided should be tested first were the austenitic stainless steels

with emphasis on those with nitrogen strengthening, a few aluminum alloys, and the fiberglassepoxy Industrial laminates. The behavior of

welds made in the stainless steels also appeared

to be a matter of concern.

4.1 Structural alloys

Considerations of strength, availability, and

the requirement for a nonmagnetic alloy with an expansion behavior close to that of the conductor led us to consider the austenitic stainless

steels as the primary alloys for low temperature

service. The nitrogen strengthened versions are particularly appealing because of their increased low temperature yield strength. For example

the yield strength at 76K of AISI 304 (a Cr-Nl

alloy) is Increased by a factor of three by the nitrogen additions which give 304N. The ultimate strength remains about the same. Other

nitrogen strengthened stainless steels of the composition Cr-Nl-Mn-N, such as the 21-6-9 alloy,

have also been Investigated. They have somewhat

higher yield strengths. Our measurements on these steels Indicate a decreased, but satisfactory, toughness at 4K. The low temperature fatigue crack growth behavior of 21-6-9 and 316 LN

is satisfactory, but recent data indicate a significant acceleration of crack growth in 304 N. Data comparing it to 304 are given In Figure 2 from [5j. This much of an Increase In crack

growth could be a serious problem in cyclic

operation. The 21-6-9 alloy has a good combination of properties, but a matching filler metal

for welding Is needed. To date, 316 LN appears

to be the most appealing structural alloy except

for the minor matter of price, but again we need more measurements to make a reasonable evaluat ion

Aluminum alloys are good candidates for struc- tural applications at low temperatures. Their

strength Is close to that of stainless steel,

but welded sections require strengthening heat

treatments which are impractical in most applications. However, bolted construction may be possible. The alloys tend to have low modulll

and high thermal expansion, both of which can occasionally present problems. We have measured

the fracture toughness and crack growth behavior

of 2219-T87 alloy at 4K and found It to be ac- ceptable in all respects.

4.2 Welds

One of our goals for this project, obviously,

is to find an alloy which not only has acceptable base metal properties, but also can be
Ing the mechanical, thermal, and electrical properties of GIOCR and GllCR samples from several

manufacturers. The laminates are anisotropic In most properties because of the layup method of

manufacture, so that a number of measurements are reaulred to characterize them completely.

Typical compressive strength data are shown In Figure 4 and the behavior of the thermal conductivity In Figure 5. Both figures are from the NBS

report on this program described earlier. In general, cooling to cryogenic temperatures Is beneficial to all of the properties of Interest. We have not yet done sufficient work to dependably evaluate the variability between manufacturers, but the preliminary results are promising.

CME

3

b

C3

QC

h“

C/3 UJ>

C/>

(/) UJ

OC

Q.O

CJ

Figure 4. Temperature dependence of the ultimate compressive strength for industrial fiberglass epoxy laminates.

epoxy laminates as a function of tem- perature.

5. STRESS AND FATIGUE IN CONDUCTORS

Whenever real structures are subjected to stress, they expand or contract to accommodate

It. In large superconducting magnets these

strains can seriously affect the superconducting

properties of the current-carrying conductors.

This degradation is especially serious for the

high field superconductors, which are brittle

materials, and It is this parameter which ultimately determines the amount of structure required for the magnet. A number of papers on this

topic were presented at the latest International

Cryogenic Materials Conference [7]. Figure 6,

from the review paper by Ekin, summarizes the

effect of strain on the superconducting critical

current of nlobium-tln conductors. Clearly, a relatively small strain can cause a significant

decrease in the current.

ik* ^ 09-

1* 1

s ^ * 08- •

*

Tc , * /cm * *0 07-

Rupp * A • *

Ekin o A Easton 8 S :hwal| 0.6- A

1 ^ ^ J 1 1 1 1 L_J 1 1 1 Lr -10 -08 -06 -04 -02 0 02 04 06 08 10x10'^

Intrinsic Strain €q Figure 6. Critical current as a function of intrinsic strain In nlobium-tln superconductor. Current values have been

normalized to the maximum in each

case

A further difficulty arises when conductors

are subject to fatigue, even at relatively low

stress levels. The pure copper stabilizer described earlier requries a low electrical resistivity in order to properly perform its function

of protecting the superconductor. Fatigue can cause the resistivity to increase significantly.

Figure 7, from the paper just mentioned, shows

this effect for copper on a nloblum-titanlum

superconductor. This ductile superconducting

alloy Is not seriously affected by the strain

levels shown, but the reslsltlvlty of the protective copper Increases very rapidly with

cycling, and it may well be this effect which

will determine the allowable strain limits for

the conductor composite In many applications.

422
Figure 2. Low temperature crack growth rate data comparing nitrogen strengtened

AISI 304 stainless steel with the

standard alloy.

Joined successfully. This means that the joint

should have properties approximating those of

the base metal. While this condition occurs frequently at room temperature, significant differences arise as the temperature is lowered.

Our measurements have sho%m, for example, that

the toughness of stainless steel welds Is marginal at 4 K In all cases Investigated to date.

We have been Investigating the metallurgical

factors which cause this reduction In toughness.

We have found that one of the most deleterious

factors Is the ferrite content of the welds as Indicated In Figure 3 for 316 L [6]. The ferrite not only degrades toughness but makes the

welds ferromagnetic. Sensitization (grain boundary precipitation of carbides) also seriously

degrades the toughness. Part of our program has

been to attempt to develop a new filler metal

for welding the 21-6-9 alloy which has very nice

properties in the base metal. A number of promising candidates have been evaluated but none

has proven to be acceptable yet. We are now evaluating the effect of various welding pro- cesses on the properties of the final weld;

preliminary results show significant variations .

Figure 3. The effect of ferrite content on weld

toughness In AISI 316 stainless steel

at cryogenic temperatures.

4.3 Industrial laminates

Fiberglass epoxy Industrial laminates are ex- tensively used In magnet construction, primarily as electrical Insulation, but occasionally

as thermal standoffs and for dewars . The most common are the NEMA (National Electrical Manu- facturers Association) GIO and Gil laminates.

There are a number of problems with these materials: their compositions tend to be proprietary

and the NEMA specifications, which are primarily

electrical, do not require any specific composition which can lead to large variations in the mechanical properties; the low temperature data

base Is very limited, and because of material

variability, the existing data may not be applicable to a particular material; the effects of radiation at low temperatures are not known, but

it is thought that they may be particularly sus- ceptible to damage to the extent of limiting the magnet lifetime. Our program has addressed the first two of these concerns directly and Oak

Ridge National Laboratory is now carrying out a program to evaluate the extent of the radiation

damage probl‘»m.

In cooperation with NEMA and several manufacturers, we have reached agreement on specific

formulations and preparation procedures for mat- erials for cryogenic applications. These are termed GIOCR and GllCR. There is also a GllCR(BF) grade which usee a boron -free glass

and, thus, may be more radiation resistant.

These materials should have uniform properties

regardless of their source. We are now evaluat423
7. REFERENCES

Figure 7. Degradation of the residual resistance ratio, R(295 K)/R(9 K), of the

copper stabilizer on a coimnerclal niobium-titanium conductor. The strain

range Is Indicated on each curve.

6. CONCLUSIONS

Much remains to be learned about the properties of materials at low temperatures as they

relate to fusion magnet applications. Large,

expensive structures, especially those Involved

In base-load power supply, require a reliability

which can only be ensured by an adequate low tem- perature data base. Such a data base does not exist at present. In the realm of fusion energy, so many problems arise due to the very high

temperatures Involved that we often tend to lose

sight of the fact that there are significant problems at the other end of the temperature scale

also

[1] F. R. Flckett, Proc. Fourth Int. Conf. on Magnet Tech. (1972), USAEC CONF- 72098 n. 4^ ^

[2] F. R. Flckett, Proc. Fifth Int. Conf. on Magnet Tech . (1975). N. Sacchettl, M. Spadonl and S. Stlpclch, eds. Laboratorl

Nazlonall del CNEN, Frascati, Italy, p. 659.

[3] F. R. Flckett, Proc. Sixth Int. Cryo. Eng.

Conf . IPC Business Press, Sussex, England

(1977), p. 20.

[4] Materials at Low Temperatures , R. P. Reed

and A. F. Clark, eds. In preparation.

[5] R. L. Tobler and H. I. McHenry, Ibid Ch.

8 .

[6] R. P. Reed, H. I. McHenry and P. Stelnmeyer,

Nuclear Engineering and Design, To be published .

[7] Advances In Cryogenic Engineering , Vol. 24,

K. D. Tlmmerhaus, R. P. Reed and A. F. Clark, eds. Plenum Press, New York

(1978).
VAIL WORKSHOP

A25

VAIL WORKSHOP

R. P. Reed

Fracture and Deformation Division

National Bureau of Standards

A major aspect of the technology transfer part of the program is

the workshop. These meetings are held each year with the goal of

bringing together representatives from all parts of the materials community

who are involved with low temperature applications. The attendees

at this year's workshop, held at Vail, Colorado, October 1 6-18, 1 979,

represented just such a diverse group including representatives of

industry, national labs, MFE projects, and universities. The following

pages give the attendance list and the program for the workshop. No

proceedings are published, but extended abstracts of the presentations

were provided to the attendees.
PROGRAM

NBS/DoE Workshop on Materials at Low Temperatures

October 16-18, 1979

Vail, Colorado

Tuesday, October 16

Session 1 (9:00-12:00): Large Magnet Programs

R. P. Reed, Chairman

DoE Office of Fusion Energy - Overview E. N. C. Dalder, LI NBS Program R. P. Reed, NBS

Large Coil Project (LCP) C. J. Long, ORNL

MHD P. G. Marston, MIT

Energy Storage H. Boenig, LASL

MFTF Magnets A. J. Hodges, LLL

Session 2 (2:00-5:00): Design, Materials and Fabrication

E. N. C. Dalder, Chairman

General Dynamics' Designs for Fusion and MHD

General Electric's Design for Fusion and MHD

Westi nghouse' s Design for the LCP

Argonne's MHD Magnet Program

Airborne MHD Magnet

Pulsed Coils

Wednesday, October 17

Session 3 (9:00-12:00): Structural Alloys

D. T. Read, Chairman

Structural Alloys for Liquid Helium Service

Ferritic Alloys for 4 K Service

Stainless Steel Alloy Development

Interstitial Strengthening of Stainless Steel

Phase Transformations and Magnetic Transitions

Session 4 (2:00-5:00): Welding

H. I. McHenry, Chairman

Aluminum Welding

Stainless Steel Welding-Filler Metals

Stainless Steel Welding-Methods

Stainless Steel Welding-Properties

US/USSR Exchange Program

Thursday, October 18

Session 5 (9:00-12:00): Composites and Nonmetallics

M. B. Kasen, Chairman

Cryogenic Grade Glass Fabric Epoxy Laminate

Cryogenic Grade Glass Mat Epoxy Laminate

Composite Resin Matrix Selection

Radiation Effects

Interlaminar Shear Strength

J. L. Christian, GD

C. L, Linkinhoker, GE

T. R. Leax, W

L. Genens, ANL

R. K. Miller, GD

S. K. Singh, W

H. IMcHenry , NBS

J. W. Morris

,

Jr.

,

R. H. Espy, Armco

R. L. Tobler, NBS

H. W. King, Dalhous

R. A. Kel sey

,

A1 coa

D. J. Kotecki , McKay

D. M. Parker, GE

T. A. Whi ppl

, NBS

R. P. Reed, NBS

R. E. Schramm, NBS

J. L. Olinger, Owens Corning

R. L. Kolek, W

C. J. Long, ORNL

H. Becker, MIT

42G
LIST OF ATTENDEES

Name

Don Beard

Herb Becker

J. R. Benzinger

Dick Blaugher

J. G. Y. Chow

Jack Christian

George D. Covington

George Cypher

Ed Dalder

Mike Dew

John Elmer

R. H. Espy

Fred R. Fickett

Lyle Genens

C. Goodzeit

Don Harwell

George Heim

Jeff Hodges

J. Horvath

Jess W. Hurt

M. B. Kasen

Ron A. Kelsey

S. H. Kim

Hubert W. King

C. Klamut

Robert Kolek

D. J. Kotecki

J. Kotora

David C. Larbalestier

T. R. Leax

A. L. Li by

C. Linkinhoker

Joe Long

NBS/DoE Workshop

October 16-18, 1979

Vail , Colorado

Affil i at ion

U.S. Department of Energy

M.I.T.

Spaulding Fibre

Westinghouse

Brookhaven National Lab

General Dynamics Convair

CTI-Cryogeni cs, INCRA

Lawrence Livermore Lab

Intermagneti cs

NBS

Armco

NBS

Argonne

Brookhaven National Lab

Cryogenic Energy Company

Fermi Lab

Lawrence Livermore Lab

General Dynamics

CVI

NBS

A1 coa Argonne

Dalhousie University

Brookhaven National Lab

Westi nghouse

Tel edyne

Argonne

University of Wisconsin

Westi nghouse

Rockwel

General Electric

Oak Ridge National Lab
LIST OF ATTENDEES

(continued)

Name Affil iation

Kim W. Mahin

Peter Marston

H. I. McHenry

John Mikurak

Ron Miller

Gerald Mordhorst

J. W. Morris, Jr.

Lawrence Berkeley Lab

M.I.T.

NBS

Sandvik, Inc.

General Dynamics Convair

Cryogenic Energy Company

University of California

J. L. Olinger Owens/Corni ng

Dave Parker

Clyde V. Prestowitz

Jeffrey V. Prose

General Electric

Arcos

CVI Corporation

David T. Read

Richard P. Reed

H. Reid

Bill Roden

M. Rosen

Robert E. Rov^lands

NBS

NBS

Tel edyne

General Dynamics

Lawrence Berkeley Lab

University of Wisconsin

Ron Scanlan

Ray Schramm

Robert Schermer

Oscar W. Seth

John Shepic

S. K. Singh

F. Spinella

E. R. Szumachowski

Arcos

NBS

Los Alamos Scientific Lab

Chicago Bridge and Iron

Martin Marietta

Westinghouse

General Dynamics

Tel edyne

Ralph L. Tobler NBS

T. A. Whipple

Dick Wolgast

NBS

Lawrence Berkeley Lab

430


PROGRAM DESCRIPTION

The overall objective of the program is to assist in the design,

construction, and safe operation of low temperature magnetic fusion

energy (MFE) systems, especially superconducting magnets, through effective

materials research and materials technology transfer. The specific steps

taken to achieve this objective are: (1) evaluation of low temperature

materials research needs specific to MFE devices; (2) development and

monitoring of a research program to acquire the necessary data; and

(3) insuring rapid dissemination of the data to potential users through

personal contacts, publications, and workshops.

Accomplishment of the first specific objective started with the

publication of the Survey of Low Temperature Materials for Magnetic

Fusion Energy in March 1977. By continuing interactions with all low

temperature design, construction, and measurement programs, such as the

Large Coil Project, we keep in touch with new problems as they arise.

This year's contribution to accomplishment of the second objective is

described here in Table 1 in the form of an outline of the research projects. The results are contained later in this report. The third objective is satisfied in part by these annual reports and by the series

of NBS-DoE Workshops on Materials at Low Temperatures, which are held each

fal 1 in Vail , Colorado.

433

Table 1. Outline of the NBS/DoE program on materials studies for

magnetic fusion energy applications at low temperatures.

Program Area Organi zation Program Description

A. Alloys for Low Temperature Service

1. Characterization NBS Fracture toughness, fatigue crack

growth rates, and tensile tests at

295, 76, and 4 K on AISI 304, 316,

304LN, and 316LN.

Tensile, fracture toughness, and

fatigue crack growth rate tests

conducted on three carbon concen- trations, each with three nitrogen

concentrations of 304 base alloy.

Tensile tests of AISI 316L conducted

at some 20 different temperatures

between 4 and 295 K. Empirical formulation of temperature dependence of

yield strength, using high temperature ORNL data.

B. Physical Properties of Stainless Steels

1. Elastic properties NBS Measurement of the elastic properties

of AISI 310 and 316 and nitrogenstrengthened stainless steels to determine temperature dependence to 4 K.

3. Effects of temperature NBS

on strength

2. Effects of C and N NBS

on strength and

toughness

C. Welding

1. Toughness NBS Evaluate the toughness at 4 K of

austenitic stainless steel welds

for LCP and MFTF and of experimental

welds supplied by filler metal producers. Relate toughness values to

metallurgical character! sties , such

as grain size, nitrogen content, and

ferri te

2. Aluminum welding Alcoa Determine the tensile and notch

tensile properties at room temperature and 4 K of welds in 5083-0

plate. GMA welds using 5183 electrode

in 51 mm thick plate were provided

for evaluation by six aluminum

producers and fabricators.
3. Mul ti -process study Tel edyne-McKay

D

1. Characterization

1 . NBS/DoE Workshop

A series of test weldments were prepared using three welding processes: gas tungsten-arc, gas

metal-arc and submerged arc. All

welds were in 25-mm thick 304L

stainless steel with 316L filler

wire. Tests were conducted to determine room temperature tensile

properties, 76 K Charpy toughness,

and microfi ssuring susceptibility.

Nonmetal lies for Magnet Structures

NBS Analyze the sources of cryogenic

property variability in G-IOCR and

G-llCR industrial glass-epoxy

laminates including a correlation

of fracture mode with mechanical

properties under tension, compression, and shear. Cooperate with

national laboratories in assessing

radiation resistance of the CR

grades including variants containing boron-free E glass. Establish

the lower bound of cryogenic mechanical properties for cottonphenolic industrial laminates. Review literature to establish state

of knowledge of polymer film properties at cryogenic temperatures.

E. Technology Transfer

NBS An annual workshop to present re- search results to the fusion com- munity, to discuss new problems, and

to promote interaction between interested parties.

436
ORGANIZATIONAL CONTACTS

Listed here are the people responsible for the various major aspects

of the program so that those with specific technical questions may contact

the individual directly. Late in the contract year responsibility for program monitoring was transferred to D. Beard after E, Oalder left OFE to

work at Lawrence Livermore Laboratories.

Department of Energy, Office of Fusion Energy, Washington, D.C. 20545

Program Monitor D. Beard

National Bureau of Standards, Boulder, Colorado 80303

Program Manager R. P. Reed

Welding

Elastic Properties

Nonmetal 1 ics

Structural Alloys

H. I. McHenry, T. A. Whipple

H. M. Ledbetter

M. B. Kasen

R. P. Reed, R. L. Tobler

(301 ) 353-4958

(FTS) 233-4958

(303)499-1000, X3870

(FTS) 323-3870

X3268

X3443

X3558

X3421

Mechanical Properties Data Center, 13919 West Bay Shore Drive, Traverse City,

Michigan 49684

Program Manager R. C. Braden

Teledyne McKay, 850 Grantley Road, York, Pennsylvania 17405

Program Manager D. Kotecki

Aluminum Company of America, Alcoa Center, Pennsylvania 15069

Program Manager R. A. Kelsey

(616) 947-4500

(71 7) 845-7581

(412)339-6651 ,X2657

437
NBS-114A (REV. B-7S)

U.S. DEPT. OF COMM.

BIBLIOGRAPHIC DATA

SHEET

1. PUBLICATION OR REPORT NO. NBSIR 80-1627

4. TITLE AND SUBTITLE

Materials Studies for Magnetic Fusion Energy Applications at

Low Temperatures - III

5. Publication Date June 1980

ganization Code

7. AUTHOR(S)

R. P. Reed, Editor

8. Performing Organ. Report No.

9. PERFORMING ORGANIZATION NAME AND ADDRESS

NATIONAL BUREAU OF STANDARDS

DEPARTMENT OF COMMERCE

WASHINGTON, DC 20234

Project/Task/Work Itelt Ha,

11. Contract/Grant No.

12. SPONSORING ORGANIZATION NAME AND COMPLETE ADDRESS rsfreer. City, state,, ziP) 13. Type of Report & Period Covered Department of Energy

Office of Fusion Energy

Washington, DC 20545 14. Sponsoring Agency Code

15. SUPPLEMENTARY NOTES

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

16. ABSTRACT (A 200^word or leaa factual aummary of moat aignificant information. If document includes a significant bibliography or

literature aurvey, mention it here.) The reports presented here summarize the work of the low temperature materials

research project for the third year of the program. The various projects are outlined and the research results are presented. The major thrust of the measurements

has been the evaluation of the low temperature mechanical and physical properties of

stainless steel base metal and welds, with particular emphasis on the nitrogenstrengthened stainless steels. Aluminum alloys have received some consideration

also. Work has also been done on the production and standardization of nonmetallics

primarily industrial laminates, for low temperature applications and on the measurement of their properties at cryogenic temperatures. The third NBS/DOE Vail workshop

was held in October 1979. A brief description is given of that program.

17. KEY WORDS ('six to twelve entries; alphabetical order; capitalize only the first tetter of the first key word unless a proper name;

aeparated by aemicotona)

Composites; low temperature; mechanical and physical properties; structural alloys;

welding

18. AVAILABILITY [^Unlimited

I I For Official Distribution. Do Not Release to NTIS

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