NBSIR 75-828

:

I

SEMI-ANNUAL REPORT ON MATERIALS RESEARCH

IN SUPPORT OF SUPERCONDUCTING MACHINERY

R. P. Reed, J. G. Hust, M. B. Kasen, H. M. Ledbetter, D.T. Read,

E.R. Naimon, R. E. Schramm, L. L. Sparks, R. L. Tobler,

and W. F. Weston

8 WW

pill

Cryogenics Division

Institute for Basic Standards

National Bureau of Standards

Boulder, Colorado 80302

January 1976

Prepared for

Advanced Research Projects Agency

1400 Wilson Boulevard

Arlington, Virginia 22209

NBSIR 75-828

SEMI-ANNUAL REPORT ON MATERIALS RESEARCH

IN SUPPORT OF SUPERCONDUCTING MACHINERY

R. P. Reed, J. G. Hust, M. B. Kasen, H. M. Ledbetter, D.T. Read,

E.R. Naimon, R. E. Schramm, L. L. Sparks, R. L. Tobler,

and W. F. Weston

Cryogenics Division

Institute for Basic Sandards

National Bureau of Standards

Boulder, Colorado 80302

January 1976

Prepared for

Advanced Research Projects Agency

1400 Wilson Boulevard

Arlington, Virginia 22209

U.S. DEPARTMENT OF COMMERCE, Rogers C. B. Morton, Secretary

James A. Baker, III, Under Secretary

Dr. Betsy Ancker-Johnson , Assistant Secretary for Science and Technology

NATIONAL BUREAU OF STANDARDS, Ernest Ambler, Acting Director

SEMI-ANNUAL REPORT ON MATERIALS RESEARCH

IN SUPPORT OF SUPERCONDUCTING MACHINERY

Sponsored by

Advanced Research Projects Agency

ARPA Order No. 2569

Program Code 4D10

August 10, 1974 - August 9, 1975

Program Director

Dr. E. C. van Reuth

Materials Sciences

Advanced Research Projects Agency

1400 Wilson Boulevard

Arlington, Virginia 22209

Program Manager

Dr. R. P. Reed

Cryogenics Division

Institute for Basic Standards

National Bureau of Standards

Boulder, Colorado 80302

The views and conclusions contained in this document are those of the authors and should

not be interpreted as necessarily representing the official policies, either expressed or implied, of the Advanced Research Projects Agency or of the U.S. Government.

iii
Abstract

Results are reported of a six-month study, March through August 1975, on candidate

materials for superconducting machinery. The results cover five areas—advanced .composites,

elastic properties, fatigue resistance and fracture toughness, magneto thermal conductivity,

and thermal conductivity. Material properties were studied over the temperature range 4 to 300 K. Materials studied include: oxygen-free copper; copper-nickel alloys; a precipitation-hardening copper alloy; invar; nickel-chromium-iron alloys; stainless steels; and the

composite materials boron/aluminum, boron/epoxy, S-glass/epoxy , graphite/epoxy , and an organic-f iber/epoxy. Some notable results of the study are: the first 4 K fatigue data

on a composite material; a ten-fold increase in the fatigue life of a uniaxial glass/epoxy

composite between room temperature and liquid-hleium temperature; the first 4 K fatigue

fracture toughness studies on a nitrogen-strengthened chromium-nickel-manganese steel,

which show this material has higher yield strength and adequate toughness compared to conventional stainless steels; room-temperature elastic properties of a copper-cadmiumchromium precipitation-hardening alloy, which are quite different from those of unalloyed

copper and show a non-parallel behavior of the shear modulus and the bulk modulus; the

thermal conductivity of 304 stainless steel may be reduced one third at 4 K by a 6 MA/m

(80 kOe) magnetic field; the first systematic study of the tensile properties of fiberreinforced composite materials between room temperature and liquid-helium temperature.

This work was supported by the Advanced Research Projects Agency of the U.S. Department

of Defense.

Keywords: Composites; copper alloys; elastic properties; engineering materials;

fatigue; fracture; iron alloys; cryogenic temperatures; mechanical properties; nickel

alloys; superconducting machinery; thermal conductivity.

iv
Report Contents

Page

ADVANCED COMPOSITES

R. E. Schramm and M. B. Kasen 1

Reprint: Mechanical and Thermal Properties of Filamentary-Reinforced

Structural Composites at Cryogenic Temperatures. 1: GlassReinforced Composites; M. B. Kasen; Cryogenics 15 (1975) 327-49 22

Reprint: Properties of Filamentary-Reinforced Composites at Cryogenic

Temperatures; Composite Reliability , ASTM STP 580, Am. Soc. Testing

Mater., Philadelphia (1975) 45

Preprint: Static Tensile Properties of Boron-Aluminum and BoronEpoxy Composites at Cryogenic Temperatures; R. E. Schramm and M. B. Kasen; Proc. ICMC-CEC, Kingston, Ontario, (July 1975), forthcoming 58

ELASTIC PROPERTIES OF ENGINEERING MATERIALS AT CRYOGENIC TEMPERATURES

H. M. Ledbetter, E. R. Naimon, and W. F. Weston 71

Low-Temperature Elastic Properties of some Copper-Nickel Alloys

H. M. Ledbetter and W. F. Weston 73

Elastic Properties of a Copper-Cadmium-Chromium Precipitation-Hardened

Alloy

H. M. Ledbetter 77

Low-Temperature Elastic Properties of Invar

H. M. Ledbetter, E. R. Naimon, and W. F. Weston 81

Reprint: Dynamic Low-Temperature Elastic Properties of Two Austenitic

Nickel-Chromium-Iron Alloys; W. F. Weston, H. M. Ledbetter, and E. R. Naimon; Mater. Sci. Eng. 20 (1975) 185-94 95

Reprint: Low-Temperature Elastic Properties of Four Austenitic Stainless Steels; H. M. Ledbetter, W. F. Weston, and E. R. Naimon; J. Appl.

Phys. 46 (1975) 3855-60 105

FATIGUE RESISTANCE AND FRACTURE TOUGHNESS OF ENGINEERING MATERIALS AT CRYOGENIC TEMPERATURES

R. L. Tobler, D. T. Read, and R. P. Reed Ill

Fatigue Resistance of a Uniaxial S-Glass/Epoxy Composite at Room and

Liquid Helium Temperatures

R. L. Tobler and D. T. Read 113

Tensile and Fracture Behavior of a Nitrogen-Strengthened, ChromiumNickel-Manganese Stainless Steel at Cryogenic Temperatures

R. P. Reed and R. L. Tobler 131

MAGNETOTHERMAL CONDUCTIVITY

L. L. Sparks 150

Preprint: Magnetothermal Conductivity of Selected Pure Metals and

Alloys; L. L. Sparks; Proc. ICMC-CEC, Kingston, Ontario (July 1975),

forthcoming ' 160

THERMAL CONDUCTIVITY

J. H. Hust 171

v
Disclaimer

Tradenames of equipment and materials are used in this report for clarity and to conform with standard usage in the scientific and engineering literature. Selection of

materials for discussion and examination with regard to application in superconducting

machinery is based on properties reported in the literature, and must be regarded as preliminary and tentative. In no case does such selection imply recommendation or endorsement by the National Bureau of Standards, nor does it imply that the material

or equipment is necessarily the best available for the purpose.

vi
NBSIR

SEMI-ANNUAL REPORT ON MATERIALS RESEARCH IN

SUPPORT OF SUPERCONDUCTING MACHINERY

ADVANCED COMPOSITES

R. E. Schramm and M. B. Kasen

Cryogenics Division

Institute for Basic Standards

National Bureau of Standards

Boulder, Colorado 80302

October 1975
Summary: Advanced Composites

Phase II of the experimental program is now in progress. This report contains 295 K,

76 K, and 4 K tensile measurements of elastic, plastic, and ultimate tensile properties

necessary to characterize five fiber-reinforced composites selected for their potential use

at 4 K: 5.6 mil boron-6061 aluminum; 5.6 mil boron-5505 epoxy; S-901 glass NASA resin 2; type A graphite-NASA resin 2; and Kevlar 49*-NASA resin 2. Compression tests are in progress.

*The use of trade names in this paper in no way implies endorsement or approval by NBS

and is included only to define the material.

2
Contents: Advanced Composites

Page

1.0 Review 4

2.0 Phase II: Characterization of the Mechanical Properties of

Uniaxial Composites at Cryogenic Temperatures 4

2.1 Introduction 4

2.2 Materials 4 2.3 Procedures 5 2.4 Results and Discussion 6 2.5 Future Work 7

3.0 References 7

4.0 List of Figures 8

5.0 List of Tables 10

Appendix I: Symbols 21

Appendix II: M. B. Kasen, "Mechanical and thermal properties of

filamentary-reinforced structural composites at cryogenic temperatures. 1: Glass-reinforced composites,"

Cryogenics 15 (6), pp. 327-349 (1975) 22

Appendix III:M. 3. Kasen, "Properties of Filamentary-Reinforced

Composites at Cryogenic Temperatures," Composite

Reliability , ASTM STP 580, American Society for

Testing and Materials, 1975 45

Appendix IV: R. E. Schramm and M. B. Kasen, "Static Tensile

Properties of Boron-Aluminum and Boron-Epoxy

Composites at Cryogenic Temperatures," Paper T3,

Cryogenic Engineering Conference, Kingston,

Ontario, July 22-25, 1975 (to be published) 58

3
1.0 Review

Previous ARPA reports [1,2] included two extensive literature reviews on the cryogenic

properties of glass-filament and advanced-fiber-reinforced structural composites. One has

recently appeared in Cryogenics [3] and the other is now in-press for a Special Technical

Publication of the American Society for Testing and Materials [4]. A reprint and preprint

of these papers are Appendices II and III. This information clearly indicated the almost

complete lack of 4 K data directly applicable to the design of superconducting machinery.

It was possible, however, to select five candidate composites for structural use and we

set out to obtain fundamental static tensile and compressive properties to make it possible

to compute design values for complex, low-temperature structures by macromechanical analyses.

Commercial suppliers began fabricating the following composites:

5.6 mil Boron-6061 Aluminum

5.6 mil Boron-5505 Epoxy

S-901 Glass-NASA Resin 2 Type A Graphite-NASA Resin 2 Keylar 49-NASA Resin 2.

While awaiting material delivery, Phase I preliminary tests were begun to determine

working procedures. For these tests, two commercial suppliers provided boron-epoxy and

graphite-epoxy materials. With these, we developed a grip system that yielded a high fraction

of satisfactory fractures and also determined that cryogenic environments changed most mechanical properties relatively little from the room temperature values for these composites.

With this background, the boron-aluminum became the first material available for our Phase II characterization program. Both this preliminary data and the boron-aluminum data appeared

in the last ARPA report [5]. Two previously unseen phenomena did appear among the results:

a two-part elastic region in the graphite-epoxy longitudinal (0°) specimens and a serrated

stress-strain curve (discontinuous yielding) in the boron-aluminum transverse (90°) and

crossply (+ 45°) specimens at 295 K and 4 K but not at 76 K.

The boron-aluminum measurements and those subsequently made on boron-epoxy were recently

presented in a talk at the International Cryogenic Materials Conference in Kingston, Ontario,

Canada (July 22-25, 1975). The paper that will appear in the proceedings of this meeting f6]

is Appendix IV.

2.0 Phase II: Characterization of the Mechanical Properties

of Uniaxial Composites at Cryogenic Temperatures

2.1 Introduction

This program constitutes the first attempt to characterize fiber-reinforced composites

at cryogenic temperatures and obtain the basic laminar properties needed for macromechanical

analysis. The present report deals with static tensile properties; compressive tests are now in progress. Other NBS investigators under ARPA contract are making plans for dynamic

mechanical tests (H.M. Ledbetter) as well as thermal measurements (J.G. Hust)

.

While the boron-aluminum data appeared in the last report, it will be included here

again to facilitate comparison and for completeness.

2.2 Materials

Each of the five composites came from the suppliers as trimmed, flat sheets that were

15.2 x 55.9 cm (6 x 22 in) for the boron-aluminum, and 27.9 cm (11 in) square for all the

polymer-matrix materials. Uniaxial layups were six plies thick for the longitudinal tests

and 15 plies for the transverse tests; crossply plates were ten plies thick. Table 1 gives the nominal thicknesses of the various sheets.

The suppliers stated that the void content should be < 2 volume % and that NaOH

extraction on two specimens of boron-aluminum indicated an average of 47.2 volume % boron.

Our own radiography of all plates revealed no visible voids. The tungsten cores of the boron

filaments were visible in the radiographs, thereby permitting a spot check of their axial

4
alignment in the finished specimens. In all cases, the filament alignment was within 1/2°

of nominal. The boron-aluminum was tested in the F temper (as fabricated) . Processing

procedures and a fuller characterization of all composites are forthcoming from the manufacturers.

2.3 Procedures

As noted in our last report [5], all test specimens were straight-sided coupons with

no tabs. The boron-aluminum specimen length was 15.2 cm (6 in) while all other specimens were 27.9 cm (11 in) long. Specimen widths were 1.3 cm (0.5 in) for longitudinal tests

and 2.5 cm (1 in) for transverse and crossply tests. Cutting was done by a diamond circular saw lubricated and cooled by a stream of water soluble oil except for the Kevlar 49 composite

for which a dry slitting saw was necessary. Light hand buffing with emery paper further

minimized any imperfections on the already very smooth edges. The Kevlar 49 material was

again an exception since this process caused fuzzy edges.

Two commercial strain gages cemented directly to the surface monitored elongation and

Poisson's ratio in the longitudinal and crossply specimens. There was only one gage on the transverse specimens since the Poisson's ratio (^j) Gan ^ e calcul ated from the trans- verse elastic modulus (^2^ anc^ t^ie l°n§i tudir- a l values of elastic modulus (E 7 ^) and

Poisson's ratio (v 19 ) by:

V

21

V

12 E22

/E

ll* Curing the strain gage cement overnight at 345 K caused softening of the NASA resin 2; thus

it was necessary to support these specimens on a flat plate during this process.

Following ASTM D 3039-71T, Method of Test of Tensile Properties of Oriented-Fiber

Composites , pre-test conditioning of polymer-based materials was exposure for at least two days to room temperature air at 40-60% relative humidity. The grip system developed

earlier \2] was attached in a separate alignment fixture just prior to the test. Gage

ler.sohs vera 15,2 cz 6 in zzz the Long 5:ici"e-s a:.: 12,2 cm - in) for those of boronaluminum. Tests were conducted in air at 295 K, in liquid nitrogen at 76 K, or in liquid

helium at 4 K at a strain rate of 0.01-0.02 min 1

.

Outputs of the load cell and strain gages were recorded continuously during the test.

To check within-specimen repeatability, we cycled the load to about one third of its

ultimate value (within the elastic limit) and back to no load about three times before

loading to ultimate for most of the polymer-matrix specimens. In some cases, the loading

was monotonic to failure: when the ultimate load was very small (e.g., graphite-NASA resin

2 transverse, and all Kevlar 49-NASA resin 2), and for the boron-aluminum which began to yield plastically early in its loading history.

As before, a "good" or acceptable fracture was one which occurred at least one specimen

thickness away from the grips. The ultimate values from the unacceptable breaks (inside or too near the grips) were usually quite comparable to those of the good tests but they were

not included in computing averages. This criterion of determining fracture location (at

least its origin) became very difficult, usually impossible, to apply to the polymer-based

lcr.gitudir.al sreci-e-.s 2ue tc theix Eracture mode (see 2,- ?.6 su2os ar.d Discuss ion 1 .

The primary purpose of testing the crossply laminates here was to determine the inplane shear modulus (G^) • With the Petit technique [7], the calculation requires knowledge

of the uniaxial elastic properties. With the Sims and Halpin approach

[8], a shear stress-shear strain curve is constructed solely from tensile data. Because

of this independence from prior test data, we've chosen to calculate with this latter

method. The values calculated by the Petit technique and listed in the last report f5]

for boron-aluminum agree quite well with those listed here.

5
2.4 Results and Discussion

Numerical results from the individual tests are presented in Tables 2-6; the average

values, standard deviations, and coefficients of variation are in Tables 6-10. Table 11

gives the calculated shear moduli.

Discontinuities on the stress-strain diagram of boron-aluminum in the transverse (90°)

and crossply (+ 45°) orientations were graphed and discussed in our last report [5]. We

still have no explanation and have found this behavior is unfamiliar to other workers.

The two-part elastic region noted in graphite-epoxy [5] was again in evidence for the

graphite-NASA resin 2. With the load along the fiber direction, the initial elastic response

is followed by a second of higher modulus and higher Poisson's ratio; these are noted as primary and secondary values in Tables 5 and 10. The uncertainties in Poisson's ratio make

possible the argument that there is no real difference, but the change in modulus is

repeatable and usually quite distinct in the load-strain curves recorded. Again, it would

appear that this increase in stiffness could be due to straightening of the graphite

filaments.

In almost all instances, a minimum of three tests resulting in "good" fractures were

conducted at each temperature. In the case of the Kevlar 49-NASA resin 2 composite, the

initial results at 295 K and 76 K were very poor and showed that it was not worth running

the full series or any 4 K tests. These low strengths and the very low interlaminar shear

strength, causing delamination problems in handling, are signs of severe incompatability

between this particular fiber and resin. Component compatability is obviously important

and, before any extensive commitment, should be checked, e.g., with short beam shear

tests.

Figure 1 shows typical specimens fractured at 4 K. The longitudinal boron-aluminum

specimens appear similar to most normal metals while longitudinal boron-epoxy specimens

fracture into discrete pieces throughout the gage length. All of the longitudinal NASA

resin 2-matrix materials tend to explode between grips and produce a fuzzy, filamentary

mass. All transverse fractures are defined by fiber direction and are typically brittle.

There is some necking and plastic yielding in all crossply specimens; in boron-aluminum,

the specimens maintain integrity while all others fail by delamination. Figure 1 shows

only 4 K fractures (except for the Kevlar 49-NASA resin 2) , but tests at other temperatures

produced breaks of basically the same type.

Mechanical anisotropy is an outstanding characteristic of composites and all properties

must be correlated with fiber direction. The high modulus of boron marks both its composites

because of their stiffnesses in the longitudinal orientation (200-235 x 10 9 N/m2 or 29-34 x 10 6 psi) . Matrix properties dominate the transverse and crossply specimens, however, and

here the aluminum matrix shows a higher modulus by a factor of three or better over its

polymer counterparts. The longitudinal moduli are very close to being temperature independent

(< 5% change between 295 K and 4 K) while the matrix-dependent orientations show some small effects: polymer bases show some increase with lowering temperature while the

boron-aluminum generally shows a small decrease.

In most metals, the Poisson's ratio is generally about 0.3; this is roughly true for

boron-aluminum in all orientations and the longitudinal specimens of the other composites.

.For the polymer matrices, the transverse values are very low while the crossply values are much higher.

Longitudinal ultimate strengths do show some temperature dependence — increasing at cryogenic temperatures particularly for boron-aluminum (16.4 x 10 8 N/m2 or 238 x 10 3 psi)

and glass-NASA resin 2 (19.6 x 10 8 N/m2 or 284 x 10 3 psi). In the other two orientations

the metal matrix again shows a superiority.

Those composite layups in this group which do display any yielding do not have a sharp transition between elastic and plastic zones. Consequently, the determination of a proportional limit is quite difficult and hence the rather large variation in values as seen in the standard deviations. The 0.2% yield strength is a considerably less variable measure of plastic yield.

6
The polymer-matrix intralaminar shear moduli in Table 11 appear to agree reasonably

with other room temperature determinations [9]; the boron-aluminum values are, however,

quite low. This difference is probably due to the non-linear stress-strain response in

the ductile matrix invalidating assumptions necessary for the indirect + 45° method of

calculating f 10] . The boron-aluminum values listed here should serve only as lower

boundaries.

Composite properties are strongly dependent on fiber volume fraction and void content

but, where general comparison is possible with the data reviewed in Appendix III, there is

reasonable agreement with the present data with a few exceptions. In the glass-NASA resin

2 we did not see the large increase in at 4 K that would be indicated by earlier 20 K

measurements [3]. The type A graphite-NASA resin 2 here had a very low transverse tensile

strength; this may indicate low filament-resin compatability . All the Kevlar 49-NASA

resin 2 properties indicate the severe compatability mentioned above.

2.5 Future Work

Compression tests are in progress. We designed and built a special fixture and

have now seen it through several modifications. After working with a few specimen configurations, it now appears that the major hurdles have been cleared and good compression

data are possible.

Fatigue testing will begin later this fall and, for this purpose, we now have in

hand boron-aluminum and boron-epoxy materials in a (0° +45°) configuration.

3.0 References

1. Kasen, M. B. , "Advanced Composites," Semi-Annual Technical Reports on Materials

Research in Support of Superconducting Machinery , (1 Sept., 1973 to 1 March, 1974)

National Bureau of Standards - Advanced Research Projects Agency.

2. Kasen, M. B. and Schramm, R. E., "Advanced Composites," Semi-Annual Technical Reports

on Materials Research in Support of Superconducting Machinery - II , (1 March, 1974 to

1 Sept., 1974) National Bureau of Standards - Advanced Research Projects Agency.

3. Kasen, M. B. , "Mechanical and thermal properties of filamentary-reinforced structural

composites at cryogenic temperatures. 1: Glass-reinforced composites," Cryogenics 15(6),

327 (1975).

4. Kasen, M. B., "Properties of Filamentary-Reinforced Composites at Cryogenic Temperatures,"

Composite Reliability , ASTM STP 580, American Society for Testing and Materials, 1975

(to be published)

.

5. Schramm, R. E. and Kasen, M. B., "Advanced Composites," Semi-Annual Technical Reports

on Materials Research in Support of Superconducting Machinery-Ill (1 Sept., 1974 to

1 March, 1975) National Bureau of Standards-Advanced Research Projects Agency.

6. Schramm, R. E. and Kasen, M. B., "Static Tensile Properties of Boron-Aluminum and

Boron-Epoxy Composites at Cryogenic Temperatures," Paper T3, Cryogenic Engineering

Conference, Kingston, Ontario, July 22-25, 1975 (to be published).

7. Petit, P. H. , "A Simplified Method of Determining the In-plane Shear Stress-Strain

Response of Unidirectional Composites," Composite Materials: Testing and Design , ASTM STP 460, American Society for Testing and Materials, 1969, pp. 83-93.

8. Sims, D. F. and Halpin, J. C, "Methods for Determining the Elastic and Viscoelastic

Response of Composite Materials," Composite Materials: Testing and Design (Third

Conference) , ASTM STP 546, American Society for Testing and Materials, 1974, pp. 46-66.

~~

9. Advanced Composite Design Guide, Third Edition, Volume IV, Materials , Air Force

Materials Laboratory, Wright-Patterson Air Force Base, Ohio (1973).
10. Ramsey, J. E. , Waszczak, J. P. and Klouman, F. L., "An Investigation of the Non-Linear

Response of Metal-Matrix Composite Laminates," AIAA/ASME/SAE 16th Structures,

Structural Dynamics and Materials Conference, Denver, Colorado (May 27, 1974)1

4.0 List of Figures

1. Composite tensile specimens fractured at 4 K (except for Kevlar 49-NASA resin 2) All transverse 90° specimens broke in the same manner while all crossply (+ 45°)

polymer matrix specimens were similar to each other in appearance.

8
FIBER-REINFORCED STRUCTURAL COMPOSITES

4K TENSILE FRACTURES

5.6 nil Boron-6061 Altiminam

I Longitudinal (0

S.6 mil Boron-5505 Epoxy

S -901 Glass-NASA Resin 2

IType A Graphite-NASA Resin 2

Kevfar 49-NASA Resin 2

Type A Graphite-NASA Resin 2 #

5.6 mil Boron-6061 Aluminum

S-901 Glass-NASA Resin 2

76 K

(Transverse (90

Crossply ( 45°

i

Tot good Weeuuu *« National Bureau of Standards

'MIHlil!ililJj||hn|||i|l|||llMUlillllililiill!lilllii;iili!;h:;| Uj

Fig. 1. Composite tensile specimens fractured at 4 K (except for Kevlar 49- NASA resin 2) . All transverse (90 ) specimens broke in the same manner while all crossply (+45 ) polymer matrix specimens were

similar to each other in appearance.

9
1.

2.

3.

4.

5.

6.

9.

10.

5.0 List of Tables

Composite Plates Used for Tensile Mechanical Tests.

Tensile Properties of 5.6 mil Boron-6061 Aluminum Composites (Individual Specimens).

Tensile Properties of 5.6 mil Boron-Epoxy Composites (Individual Specimens).

Tensile Properties of S Glass-NASA Resin 2 Composites (Individual Specimens).

Tensile Properties of Type A Graphite-NASA Resin 2 Composites (Individual Specimens)

Tensile Properties of Kevlar 49-NASA Resin 2 Composites (Individual Specimens and

Averages)

.

Tensile Properties of 5.6 mil Boron-6061 Aluminum Composites (Averages of Specimens

Tested).

Tensile Properties of 5.6 mil Boron-Epoxy Composites (Averages of Specimens Tested).

Tensile Properties of S Glass-NASA Resin 2 Composites (Averages of Specimens Tested)

,

Tensile Properties of Type A Graphite-NASA Resin 2 Composites (Averages of Specimens

Tested).

11. In-Plane Shear Modulus (Averages of Three Specimens)

Table 1.

Composite Plates Used for Tensile Mechanical Tests

Composite Orientation Nominal 1 (cm)

'hickness

(in)

Plies

5.6 mil Boron 6061-Aluminum longitudinal (0°) 0.109 0.043 6

transverse (90°) 0.246 0.09715

crossply (+ 45°) 0.178 0.070 10

5.6 mil Boron-5505 Epoxy longitudinal (0°) 0.102 0.040 6

transverse (90°) 0.257 0.101 15

crossply (+ 45°) 0.170 0.067 10

S-901 Glass - NASA Resin 2 longitudinal (0°) 0.124 0.049 6

transverse (90°) 0.315 0.124 15

crossply (+ 45°) 0.196 0.077 10

Type A Graphite-NASA Resin 2 longitudinal (0°) 0.112 0.044 6

transverse (90°) 0.292 0.115 15

crossply (+ 45°) 0.196 0.077 10

Kevlar 49 - NASA Resin 2 longitudinal (0°) 0.124 0.049 6

transverse (90°) 0.356 0.140 15

crossply (+ 45°) 0.218 0.086 10

10
Table 2.

Tensile Properties of 5.6 mil Boron-6061 Aluminum Composites

(Individual Specimens)

Poisson'

s

Proportional 0 2% Yield Ultimate Ultimate

T emperature Elastic Modulus E Ratiof v Limit, ^ Strength. Strength.

_tu Elongation, e

(Kl

9 2 (10 N/m ) <io

b psl) 8 2 (10 N/m ) 3 (10 psi) (10

8 N/m2

) (10

3 psi) (10

i 2 N/rcT) do3 psi

)

(%)

Long tudinal (

Z95 195 28. 3 0. 320 12. 4

b 180

b 0.6b 295 194 28. 1 0. 280 12. 7 L 184 u

°- 6

h 295 194 28. 2 0 275 1" .2° 191

b

n (, b 295 217 31. 4 0. 320 12.5" 175 b 0.6

h

76 186 27. 0 0. 345 16. 9 24 5 0. is 76 196 28. 0. 344 15 . 8 230 0. 7 76 2 08 30. 1 0. 335 16 . 5 239 0. 8 76 196 28. 4 0. 351 It .0 b 233 b

0. 8

b

» 211 30. 6 0. 313 14. 9 21 0. 7

i 4 207 30. 0. 358 16 . 9 24 5 0. 8 4 191 27. 7 0. 343 16 . 0 231 0. B 4 206 29. 9 0. 360 16 . 4 238 0. 8

Transverse (90°

)

295 165 24. 0 0 247 0. 67 9. 7 1. 38 20 1 1. 74 25 3 0. 8 295 177 25. 7 0. 265 0. 54 7. 8 _ 1. 80 26 1 295 1 12 20. 6 0.212 0. 67 9. 7 1. 29 18 7 1. 46 21 1 0. 6

76 101 14. 6 0.175 0. 92 13. 3 1. 86 27 0 2. 29 b 33 2 b n o h

0. 8

'ib 155 22. 5 0. 268 0. 75 10. 9 1. 67 24. 3 2. 57 37 3 0. H 76 138 20. 1 0. 239 0. 65 9. 4 1. 94 28 2 2. 47 35 9 0. 8 76 1 16 16. 8 0. 201 0. 72 10. 5 1 . 66 24 1 2. 42 35 1 0. 9

4 180 26. 2 0.304 0. 50 7. 3 1. 97 28. 6 2. 79 40 5 0. 9 4 152 22. 1 0.256 0. 80 11. 6 2. 05 29. 7 2. 73 39 6 0. 8 4 118 17. 1 0. 199 0. 86 12. 5 1. 87 27. 1 2. 76 40 0 0. 9

Crossply ( * 45° )

295 1 12 16. 2 0. 27 3. 9 0. 90 13. 0 3. 62 52 4 13. 6 295 100 14. 5 0. 417 0. 68 9. 9 1. 01 14. 6 3. 92 56 9 295 120 17. 3 0. 375 0. 17 2. 4 0. 92 13. 3 3. 32 48 2 10. 6

76 61. 8 9. 0 0. 313 0. 08 1. 2 1. 52 22. 0 4. 11 59. 6 4. 6 76 71. 8 10. 5 0. 361 0. 20 2. 8 1. 44 20. 9 4. 64 67 2 8. 8 75 99. 0 14. 4 0. 325 0. 16 2. 3 1. 09 15. 8 3. 66 53. 1

4 109 15. 9 0. 405 0. 48 6. 9 1. 40 20. 3 4. 25 61. 7 3. 2

4 111 16. 0 0. 378 0. 26 3. 7 1. 20 17. 4 4. 06 58. 9 4. 1 4 84 12. 1 0. 337 0. 18 2. 7 1. 41 20. 5 4. 02 58. 4

aPoisson's ratios for transverse (90") specimens (v^j) were calculated from E^ and the average values of , and E

b These values resulted from an invalid fracture and were not included in the average value.

11
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15
Table 7.

Tensile Properties of 5.6 mil Boron-6061 Aluminum Composites

(Averages of Specimens Tested3

)

I emperature Elastic Modulus E

b Poisson's Ratio, v Proportional Limit. 0.2% Yield Strength. <-' V Number o(

9 , 2 8 2 Specimens

IKl (10 N/m I (10

6

psil CV(%) CV(%) (10

8 N/m2

) (10

3 psi) CV(%) (10 N/m ) (1 0

3 psi) CV(%)

Longi udinal (0°

)

295 2001 1 1 I 29. 0(1. 61 5. 7 0. 299(0. 024) 8. 2 4

76 1 17(1 1

1

28. 5(1. 6) 5. 6 0. 341(0. 005) 1.6 4

4 2(14(91 29. 6(1. 31 4. 3 0. 344(0. 022) 6. 3

Transverse (90 1

4

295 161 (181 23. 4(2. 6) 11.0 0. 241 (0. 027) 11.2 0. 63(0. 08) 9. 1(1. 1) 11.9 1. 34(0. 051 19. 4(0. 71 3. 6 3

76 1 2KI24I 18. 513. 51 18. 7 0. 22110. 041) 18. 6 0. 76(0. 1 1 ) 1 1. 0(1. 61 15. 1 1. 78(0. 141 25. 9(2. 0) 7. 8 4

• 4 1 50(31) 21. 8(4. 6) 20. 8 0. 253(0. 053) 20,8 0. 72(0. 19) 10. 5(2. 81 26. 8 1. 96(0. 091 28. 5(1. 31 4. 6 3

Crossply (

1 45°

)

2 '15 111(101 16. 01 1. 4) 9. 0 0. 396(0. 0211 5. 3 0. 37(0. 27) 5. 414. 0) 73. 0 0. 94(0. 061 13. 6(0. 9 6. 2 3

781 |-0.| 1 1. 3(2. 81 2 4. 8 0. 33310. 025) 7. 5 0. 15(0. 06) 2. HO. 81 40. 7 1. 35(0. 231 19. 6(3. 3 16. 9 3

4 1 III 1 1 51 14. 7(2. 21 14. 9 0. 373(0. 034) 9. 2 0. 31(0. 161 4. 4(2. 21 5 0. 1 1. 34(0. 121 19. 4(1. 7 8. 8 3

.Temperature

(K)

2'1S 76

4

295

76

Ultimate Strength, /m2

) (10

3 psi) CVC;

12. 7

16. 4(0. 6)

16. 1(0. 91

1..67I0. 181

2. 49(0. 08)

2. 76(0. 03)

3. 62(0. 30)

4. 14(0. 491

4. 11(0. 12)

Ultimate

Elongation, c

Longitudinal (0°

184

238(8) 233(12)

3. 4

5. 3

Transverse (90°

)

24. 2(2. 7)

36. 1(1. 1)

40. 0(0. 5)

10. 9

3. 1

1. 1

Cros sply (± 45 °

)

52. 5(4. 4)

60. 0(7. 1)

59. 7(1.81

8. 3 11.8

3. 0

0. 6

0. 8(0. 01

0. 8(0. 1 )

0. 7(0. 1

1

0. 8(0. 1 I

0. 9(0. 1)

12. 1(1. 51

6. 7(2. I I

4. 2(1. 01

14. 3

7. 2

6. 4

12. 4

31. 3

23. 8

Number of Specimens

Standard deviations are in parentheses. CV is coefficient of variation. Data from the specimens that broke at or in the grips are

not included in the calculations of ultimate strength or elongation. ''Poisson's ratios for transverse (90°) specimens (v^j) were calculated Eromv^, ^M p ancl E 22'

16
Table 8.

Tensile Properties of 5.6 mil Boron- Epoxy Composites

(Averages of Specimens Tested )

Temperature Elastic Modulus, E Poi s son s Ratio!" v Proportional Limit. 0. 2% Yield Strength.

ty Number of

(KI

9 , 2

(10 N/m ) (10 psil CV(%) CV(%) (10

8 N/m2

) (10

3 psi) CV(%) (10

8 N/m2

) (10

3 psi) CV(%| Specimens

Longitudinal (0°)

i 295 231(41 33. 5(0. 6) 1. 6 0. 228(0. 019) 8. 3 5 76 233151 33. 9(0. 7) 2i 1 0. 241(0. 012) 4. 8 5

4 238(61 34. 5(0. 6) 2. 3 0. 239(0. 014) 5. 9

Trans rerse (90*

)

5

295 17. 5(0. 0] 2. 54(0. 12) 5. 0 0. 017(0. 001) 6. 1 0. 344(0. 071) 4. 99(1. 02 20. 7 6 76 31. 2(2. 41 4. 53(0. 341 7. 5 0. 033(0. 002) 7. 3 4

4 35. 8(2. 31 5. 19(0. 32) 6. 3 0. 034(0, 003) 7. 8 3

Crossply (± 45°)

29 5 18. 1(1. 1 '. 2. 62(0. 171 6. 3 0. 817(0. 034) 4. 1 0. 406(0. 0131 5. 89(0. 19 3. 3 3. 956(0. 016) 13. 9(0. 3) 1. 8 3

. 76 32. 7(0. 7) 4. 72(0. 13) 2. 8 0. 671(0. 050) 7. 4 0. 411(0. 020) 5. 96(0. 29 4. 9 D. 951(0. 038) 13. 8(0. 51 3. 9 3

4 33. 3(1. 91 4. 83(0. 281 5. 8 0. 700(0. 051) 7. 2 0. 296(0. 028) 4. 29(0. 40 9. 3 3. 830(0. 016) 12. 0(0. 21 1.9 3

Temperature

(K

)

Ultimate Strength,

..8.., Z (10 psi) CV(%)

Ultimate

Elongation, c Number o( Specimens

295

76

295

76

295

76

16. 3(0. 6)

16. 8(0. 91

18. 2(0. 4)

0. 470(0. 035)

0. 490(0. 033)

0. 414(0. 0961

1. 30(0. 02)

1. 11(0. 02)

1. 01(0. 03)

Longitudinal (0°

237(8) 243(13)

263(6)

Transverse (90°

3. 3 0. 73(0. 04)

5. 4 0. 77(0 071

2. 2 0. 76(0. 01)

6. 81(0. 51)

7. 09(0. 47]

6. 01(1. 39)

7. 4

6. 7

23. 1

Crossply (± 45°)

18. 8(0. 3)

16. 1(0. 3)

14. 7(0. 4)

1. 4

1. 7

2. 8

0. 27(0. 06)

0. 15(0. 03)

0. 12(0. 04)

1. 88(0. 14)

0. 85(0. 05)

0. 83(0. 05)

5. 1

9; 3

1. 5

21. 7

17. 6

33. 6

7. 5

5. 3

6. 3

Standard deviations are in parentheses. CV is coefficient of variation. Data from the specimens that broke at or in the grips are not inc'uded

in the calculations of ultimate strength or elongation. 'Poisson's ratios for transverse (90°) specimens (v 2]

) were calculated fromv.,, £.., and E .

17
m cn cn cn m en

o

0) H

s

EH

5

c

rH rH vO

i1 a d d

>^ cn on

rH rH 00 H

ON CO CM

in m m

Number

of

Specimens <n eh cn cn m cn cn cn -3-

tu

ongation,

e

CV(%) 8.2 3.0

11.7 19.2 29.6 15.6 15.1 16.0 6.6

Ultimate

El

(%)

(0°)

2.8(0.2) 3.3(0.1) 3.0(0.4)

;90°)

0.6(0.1) 0.4(0.1) 0.4(0.1)

45°)

6.0(0.9) 1.3(0.2) 0.9(0.1)

B

3

Ltudinal 11.4 3

.

5

11.0

nsverse

I

25.3 17.0

ssply

(+

6.2

10.9 2.0

D ^

- tH X 0) u a

00C cn

01 o

u d

1

Long:

198(22) 284(10) 281(31)

Tra:

6.8(0.4)

13.6(3.4) 11.4(1.9)

Cro

20.1(1.2) 33.0(3.6) 34.1(0.6)

CO

0) *J . !~B •iH w z

.3 CO

6*S CM sO in en r- CV en o ON CM

w CO

cn \o ON CM a m en CM rH CM

CO *H 3 W5 3

TJ

45(0. 73(0. 80(0. 07(0. 20(0, .87(0 .59(0 .02(0 o o

s d

CO CO CO en CM m

u

4-1

(T) CN

rH B m rO r- rH <r !-» U ^•z o O cm CM CM H

ON en en o CO m

0T) r-m

COmd CM

CM ON

tn

<*- m

<u

Vh3

4-) m sO IDa\

CM

sO inON

CM

sO

O.s

cm cn cn

m cm m

p-) ON ON

V- 0J

ctj T)

18
19
Table 11.

In-Plane Shear Modulus

(Averages of Three Specimens)

Temperature Shear Modulus, G. _ 9 2 (10 N/m ) (10 psi) CV(%)

b Boron-Aluminum

295 39.7(4.0) 5.75(0.58) 10.0

76 27.5(5.5) 3.92(0.80) 20.4

4 32.6(5.5) 4.73(0.80) 16.9

Boron-Epoxy

295 4.72(0.44) 0.68(0.06) 9.3

76 9.19(0.38) 1.33(0.06) 4.1

4 9.28(0.19) 1.35(0.03) 2.0

Glass - NASA Resin 2 295 6.19(0.51) 0.90(0.07) 8.3

76 11.0(0.6) 1.60(0.09) 5.8

4 12.9(0.3) 1.87(0.05) 2.8

Graphite - NASA Resin 2 295 4.14(0.33) 0.600(0.047) 7.9

76 4.50(0.15) 0.652(0.022) 3.3

4 5.32(0.36) 0.772(0.053) 6.8

aThese values are calculated from the crossply (+ 45°) data using

the method of Sims and Halpin. Standard deviations are in

parentheses; CV is the coefficient of variation.

bThe boron-aluminum moduli are only lower limits to the true values.

20
Appendix I: Symbols

E^^ = longitudinal (0°) elastic modulus

= transverse (90°) elastic modulus

Ex

= crossply (+ 45°) elastic modulus

= longitudinal (0°) Poisson's ratio

v = transverse (90°) Poisson's ratio

= crossply (+ 45°) Poisson's ratio

G^2 = intralaminar (in-plane) shear modulus

pi

o = proportional limit; stress at which a stress-strain curve deviates

from elasticity, i.e., is no longer a straight line

ty

0 = 0.2% tensile yield strength; stress at the intersection of the

stress-strain curve with a straight line having the same slope

as the elastic modulus and offset by 0.2% along the strain scale

0 tu

= ultimate strength; highest stress attained on the stress-strain curve

£ tU

= ultimate elongation; highest strain attained on the stress-strain curve

21
Appendix II: M. B. Kasen, "Mechanical and thermal properties of filamentaryreinforced structural composites at cryogenic temperatures. 1: Glass-reinforced composites," Cryogenics 1£(6), pp. 327-349

(1975).

This article is an extensive review of the literature on the mechanical and thermal

properties of glass-reinforced structural composites at cryogenic temperatures. The

objective is to provide an understanding of the general magnitude of property values obtainable within the cryogenic temperature range, to provide a feel for the relative literature ranking of specific composite types with regard to a specific property, and

to impart an understanding of the temperature sensitivity of the property of interest. A bibliography and bibliography-property cross-reference is included. This is Part I of a two-part series. Part II will consider advanced composites.

Mechanical and thermal properties of filamentaryreinforced structural composites at cryogenic

temperatures

1: Glass-reinforced composites

M. B. Kasen

Nomenclature

tensile ultimate strength

initial tensile modulus

secondary tensile modulus

e

tu tensile ultimate strength

„tu tensile fatigue failure stress

a

fu flexural ultimate strength

initial flexural modulus

4 secondary flexural modulus

compressive ultimate strength

Ec compressive modulus

0

s' interlaminar shear strength

The primary impetus for structural composite development

arose from the need to obtain improved long-term mechanical properties at elevated temperatures or to reduce the cost of structures designed for ambient temperature use. Comparatively little effort has been expended on development

of composites for use at cryogenic temperatures. A notable exception has been the rather extensive body of work

sponsored by the USAF and NASA wherein a series of

glass-reinforced plastics were characterized to 20 K. The

other major field of cryogenic development has been con- cerned with composite reinforcement of pressure vessels

The author is with the Cryogenics Division, National Bureau of Standards, Institute for Basic Standards, Boulder, Colorado 80302, USA. This research was supported by the Advanced Research

Projects Agency of the Department of Defence under ARPA Order No 2569. Received 6 January 1975.

a

by bearing yield strength

a

1 impact strength

X thermal conductivity

AL/L thermal contraction CP

specific heat

Units

lb in-2 pounds per square inch

k lb in2 lb in"2 x 10 3

N nr2 Newtons per metre squared (Pascal)

J kg"1 K

_1 Joules per kilogram-Kelvin

W nr1 K

-1 Watts per metre-Kelvin

for aerospace use, largely exploiting the continuousfilament method of fabrication. To a large extent, the re- maining published data on composite properties at cryogenic temperatures reflects work on which the generation of cryogenic property data was peripheral to the main work

objective.

This relative lack of emphasis on cryogenic structural composites is perhaps understandable, as the majority of

such structural applications are presently satisfied by

readily available and well-characterized metals and alloys. In view of the extensive data base available on metals, it is probable that metals will continue to constitute the main

body of structural materials at low temperatures. Why then, should one consider composites? The answer

lies in the increasingly stringent demands made on materials

in advancing cryogenic technology, of which* superconductCRYOGENICS. JUNE 1975 327
ing machinery' may serve as an example. Undoubtedly, the

first generation of superconducting motors and generators

will be dependent almost entirely on metals technology. However, it is highly probable that succeeding generations of such equipment will capitalize on advanced composite

technology for reasons of increased reliability, reduced

weight, and increased efficiency, reflecting the higher specific strengths and moduli of advanced composites coupled

with a wider range of thermal and electrical properties than

are obtainable with any conventional metal. In particular, the high strength of the polvTneric-matrix composites combined with very low thermal conductivity will be advan- tageous in minimizing heat losses in critical components.

The technological problems associated with integration of

composites into superconducting machinery are three-fold:

(a) most designers lack a feel for the properties available with composites, (b) an adequate data base does not exist for composites at cryogenic temperatures, particularly at 4 K, and (c) most existing composites are optimized for

service at room temperature and above not for for cryogenic service. The current programme at NBS is aimed at these three problem areas. Our first effort has been to initiate an extensive review of what is known about the mechanical and thermal properties of composites at cryogenic temperatures. The objectives of the review are four-fold: (a) we wish to provide the

designer with a feel for the general magnitude of property

values which may reasonably be expected from a given cate- gory and class of composites within the cryogenic range. (We define a composite category- by the general reinforce- ment type, for example, glass-fibre or advanced fibre (graphite, boron, etc). We subdivide the category into composite classes by the general matrix type, for example, glass- polyester or graphite-epoxy. We further subdivide the class by referring to a composite type when a specific reinforce- ment/matrix combination is specified, for example, HT-S

graphite/X-904 epoxy.] (b) we wish to provide him with a

feel for the literature ranking of specific composite classes with regard to a specific property, (c) we wish to impart a

feel for whether the property of interest is likely to increase, remain unaffected, or decrease with lowering of temperature, and (d) we wish to define those areas in which further data are needed and to define the direction that future work

should take in optimizing composites for cryogenic service and for implementing their use in the construction of super- conducting machinery. For those with more specific in- terests we include an extensive Bibliography and Biblio- graphy-property cross-reference to simplify retrieval of

specific documents.

It must be emphasized that, although this survey is intended

to be comprehensive, the complex nature of the subject

makes it unavoidable that some works worthy of inclusion have been inadvertently omitted. The author will appreciate references to any additional material which bears on

the subject. It must also be emphasized that the results reported in the surveyed publications have not been exper- imentally confirmed by NBS and that the conclusions and

evaluations presented in this paper reflect those of the cited authors and do not imply approval, endorsement, or recommendation by NBS.

The discussion of properties given below does not take into consideration the effect of variations in fibre/resin ratio of

specific types of composites and test specimens, as this characteristic was not reported for all referenced works. Composite properties may be strongly influenced by this ratio. The property data discussed in this paper reflects actual

values and trends reported by the cited authors for specific composites. Controlled variations in many of the properties

are obtainable in practice by specific variation of the fibre content of the composite. The literature review covers 1960 to the present time, as it

is within this time span that almost all of the significant work was undertaken. We include only continuous-fibre

reinforced composites, as such composites are the primary

structural materials. The review excluded cryogenic insul- ations, superconductor composites, thin films, honeycomb

strucures, composite-overwrapped metal, or fibre-only properties. Filled composites are also excluded. There remains

the very large field of composites reinforced by a variety of

fibres in a variety of lay-ups in a variety of matrix materials, and it is with this body of data that this review is concerned. The wide v ariety of composite formulations and lay-ups are further complicated by lack of standard test procedures.

Furthermore, as the field is relatively new, much of the

earlier work was performed on relatively poorly characterized composites. We have attempted to cope with this complexity by dividing the review into two major sections: Part I treating glass-reinforced composites and Part II treat- ing the so-called advanced composites, for example, boron,

graphite, etc. The rationale for this separation is the dis- tinctly different use of these two composite categories in engineering practice, that is, glass-reinforced composites are used in applications where stiffness is not a design limitation, while the advanced composites are used where a high

modulus material is essential. Within each of these categories, we present the reader with

a series of graphs on which appear the average literature values of each property for each composite class from room

temperature into the deep cryogenic range. Admittedly,

presenting average data is in danger of being misleading, as each curve has associated with it a considerable scatter band. For this reason, we discuss the range of values associated with each curve, emphasizing those specific composite types

for which the highest values were reported. It is of utmost

importance, however, that the reader understands that the graphed data appearing in this review are class averages and

are therefore not to be used for engineering design purposes. The reader interested in design values is referred to the

specific literature references.

Scope of the literature survey We initially conducted a subject search using the data bases of the NBS Cryogenic Data Center (CDC), Defense Documentation Center (DDC). and the National Technical Information Service (NTIS). Additionally, a subject search was conducted through the volumes of NASA Scientific and

Technical Aerospace Reports (STAR), and the ASM-AIME

Metals Review. As the search progressed, a series of contract numbers were identified as being associated with studies of

the cryogenic properties of composites. The DDC and NASA data bases were then searched for all reports issued under such contracts. Finally, the DDC data base and that of the Smithsonian Science Information Exchange were

searched for current work in progress.

Organization of the Bibliography

The appended Bibliography contains 148 references. As the work progressed, it became apparent that a large part of the

328

23

CRYOGENICS . JUNE 1975
relevant data had been produced under a relatively tew contracts sponsored either by NASA or the USAF. Eleven such contracts have been listed at the start of the Biblio- graphy with references to the most pertinent publications

issued under each contract. Only final reports are listed, as they adequately summarize the data which also appeared in numerous interim reports on each project. Journal publications are listed, as they often provide a convenient review of

the subject matter contained in the comprehensive reports and are generally more readily available to the reader. The Bibliography also contains a general section, alphabetically arranged by author, listing relevant publications

sponsored by other contracts or by corporate in-house

funding. A separate part of the Bibliography itemizes handbooks or reviews which will be found useful references but which do not contain original data. Finally, a miscellan- eous reference section lists publications which are referenced

in the text but which do not contain relevant mechanical

or thermal property data. Wherever possible, the pertinent NASA or DDC code number is included to facilitate retrieval of specific publications. Corporate reports not so identified must be obtained from the corporate source. An extensive cross-reference relating mechanical and physical properties of specific types of composites to specific

literature references is included to simplify literature re- trieval by the reader. The latter includes separate references

to filament-wound pressure vessels in recognition of the importance of such applications to cryogenic technology.

A separate listing is also provided of reports containing information on the effect of combined cryogenic temperature and nuclear radiation.

Glass-reinforced composites

The mechanical and thermal properties of glass-polymeric

composites are summarized on Figs 1 — 16. Where available, data are presented for 295 K, 200 K, 77 K, and 20 K (4 K

data are almost non-existent). Straight lines connect

average values at each temperature. Absence of a data point

for a given temperature implies that no significant data was

found in the literature. An asterisk adjacent to the number

identifying a curve indicates that the data for that particular

composite type was minimal relative to that available for the other composite types included on the figure. When considering the mechanical property data, the reader should be aware that there exists no universally accepted

method of determining these properties for composites, al- though committees of the ASTM are working diligently on the problem of standardization. Tire data discussed in the present review were for the most part obtained in the course of comprehensive research programmes by reliable investigators who were concerned with obtaining the most

valid results possible. Nevertheless, it remains a possibility that some of the scatter in the data reported in the litera- ture, particularly for compression and interlaminar shear,

reflects differing test procedures. Where this has become

apparent, the data have been separated by test method.

Static mechanical properties

Composite tensile strength and modulus

The reason for the widespread use of glass-reinforced composites is evident from the tensile strength data presented on

Fig.l. No other type of composite can match the uniaxial

tensile strengths of 250—300 k lb in"2 provided by the glass- epoxy formulations. Even in the 0°/90° crossply lay-up, the glass-epoxy strength is almost equal to that of the advanced

composites in the uniaxial longitudinal configuration. Unfortunately, the moduli of glass-reinforced composites are quite low, as may be seen in Fig.2. It is this low modulus

of glass that has given impetus for development of the advanced composites. From Fig.l we see that the tensile strength is reduced about 50% at all temperatures in a 0°/90° configuration as compared to uniaxial, for example, from about 300 k lb in 2 to about 1 50 k lb in"2 for glass-epoxies. This is expected from

a 50% decrease in longitudinal fibre content, the crossply

fibres cqntributing nothing to the overall strength when

tested parallel to one fibre direction. Approximately another 25% decrease in strength is found upon going to woven cloth reinforcement of an epoxy

matrix, reflecting the decrease in load carrying capacity of

fibres which are slightly bent in the weaving process. Epoxy resins are most widely used as matrices for structural applications where maximum strength is required. This appears justified on the basis of the data of Fig.l wherein

the ultimate tensile strengths of the cloth-epoxy composites

are overall higher than those of other cloth-polymeric composites at all temperatures. The polyurethane, Teflon, phenyl silane, polyimide, and silicone-matrix composites

appear to have the poorest strength properties at 295 K, with phenolic, polybenzimidazole (also known as PB1 or

Imidite), and polyester-matrix composites being inter- mediate in strength. At 77 K, the polyimide, silicone, and

phenyl silane-matrix composites continue their relatively poor performance, while the polyurethane, polybenzimidazole, and Teflon-matrix composites have about equalled the phenolics and polyesters.

Temperature, K

Fig.l Ultimate tensile strength of glass-reinforced composites

1 — epoxy; 2 — polyurethane; 3 — phenolic; 4 — polyimide; 5 — polyester; 6 — silicone; 7 — phenyl silane; 8 — Teflon; 9 — polybenzimidazole

CRYOGENICS. JUNE 1975 24

329
Among the unaxial glass-epoxy composites, the highest re- ported strengths were about 175 k lb in"2 at 77 K for NOLring tests with S-HTS/660 FWand S-901/ERL 2256/ZZL

0820 composites.46,47 This work was directed toward fila- ment-wound pressure vessels, which has been the incentive

for much of the cryogenic composite development work. Other authors 3 ' 66 have reported strengths of 315—330

k lb in"2 at this temperature for S-901/E-787 and S-HTS/

Epon 828/Epon 1031 filament-wound composites. About

the same strength was reported for Hi-Stren glass filament- wound in a NASA Resin 2 matrix.4 The overall data range was 213-375 k lb iir2 at 77 K. A comment on two of these epoxy resins is in order so as to prepare the reader for proper understanding of these and

subsequent data. Resin E-787 will be referred to frequently

in this paper, as it is often associated with the highest re- ported mechanical properties. E-787 is the US Polymeric

Corporation designation for an Epon 828/Epon 1031/NMA/

BDMA formulation in proportions 50/50/90/0.55 pbw.

This resin is referred to as 58—68R by the Shell Chemical Company and is often referred to as the Polaris resin be- cause of its successful use in that missile. It is a conventional type of resin, not optimized for cryogenic service. By contrast, Resin 2, consisting of Epon 828/DSA/Empol

1040/BDMA in proportions 100/1 1 5.9/20/lpbw, is a bisphenolic. A epoxy system modified for low-temperature

flexibility be means of a long-chain anhydride and a high- molecular weight tricarboxy acid. This latter resin was

developed by Soffer and Molho under NASA sponsorship.5 Soffer and Molho compared the Resin 2 formulation to that of 58—68R (E-787) in a series of burst tests on metal-lined, S-901 glass filament-wound pressure vessels, their results showing Resin 2 to develop equal strength at room temperature and higher longitudinal filament stresses at 77 K and 20 K. Additionally, Resin 2 was found to have relatively greater ductility and toughness at cryogenic temperatures and to be highly resistant to cracking when thermal cycled

to 20 K. As the E-787 and Resin 2 epoxy formultions are

likely candidates for cryogenic applications, the present

report compares the property values obtained with each

resin whenever such data are available. A further discussion of resins for cryogenic service is presented in the concluding part of this paper. Continuing with the analysis of the data on Fig. 1 , it is found that among the 0°/90° data, the highest tensile strengths at cryogenic temperatures have been reported for S-901 /E-787 3 and S-901 /Resin 2

s with values ranging from 170-200 k lb in"2

. The overall range was 1 1 5-200 k lb in2

. Among the woven-cloth composites, the highest cryogenic

composite strengths were reported for the epoxy-matrix

composites 181 /modified Epon 8282 and 1581 /E-787, 3 having tensile strengths on the order of 125-145 k lb in-2

at 77 K tested parallel to the woof or warp. The phenolic- matrix composite data showed quite consistent strengths of 60-70 k lb in-2 at 77 K except for the work of Levin et al 71 who has reported 200 k lb in-2 at 77 K in a composite based on a butvar-phenolic adhesive. The polyester-matrix data

scatters from 50-80 k lb in"2 at 77 K except for one report of 100-105 k lb in"2 for 181 glass in Hetron 31 or Narmco

527 resin. 2 The Teflon-matrix composites ranged from

50—80 k lb in-2

, the highest value being reported for type

1 16 glass in TFE or FEP.80 The silicone-matrix data showed relatively large scatter from 25-70 k lb in"2 aj 77 K with

the highest values being reported for 181/Trevarno F- 1 3

1

2

and for 181/Narmco 513. 3 Few data were found for poly- urethane, phenyl, silane, and polyimide-matrix composites.

About 73 k lb in"2 was reported at 77 K for 181 glass rein- forced with the flexible polyurethane Adiprene L-100 and 60 k lb in-2 for the same reinforcement in the phenyl

silane Narmco 534. 2 Polyimides are relatively new matrix

materials, having been developed primarily for elevated- temperature use, particularly for stability. Krause et al 68 reports a comparatively low value of 43 k lb in 2 at 77 K

for a glass-polyimide composite. On the other hand, poly- benzimidazole-matrix composites, which are of the same

family as the polyimides, appear to develop very high

strength at 77 K and at 20 K, exceeded only by the

epoxies.

All of the above comparisons were made at 77 K because,

as seen on Fig.l, while the tensile strength in all cases in- creases between 295 K and 77 K, further cooling to 20 K

produces erratic results. This is more clearly illustrated by

Fig. 5a, which is a histogram illustrating the frequency with which the literature reports a given change in the ultimate

tensile strength of glass-epoxy composites upon cooling from room temperature to 77 K and to 20 K. The data are broken down as to lay-up type. Here it is seen that cooling of glass-epoxy composites from 295 K to 77 K can produce a strength increase of from 10-140 k lb in"2 with

a high probability of an increase on the order of 30—60

k lb in"2

, essentially independent of the type of lay-up. However, on cooling further to 20 K, the probability is for

a slight decrease in strength for cloth and 0°/90° crossply reinforced epoxies and a reasonably high probability that uniaxial glass-epoxies will suffer a strength degradation

which may be as high as 80 k lb in"2

. The phenolic,

polyester, phenyl silane, and polyurethane-matrix composites all showed a similar erratic behaviour. An exception appeared to be the silicone-matrix composites which

showed consistent moderate increases in strength at 20 K. The behaviour at 200 K offers few surprises except for the 181/Adiprene L-100 data 2 which indicate that the

strength of this flexible polyurethane composite rapidly

increases as temperature is lowered.

As with the tensile strength, the initial tensile modulus (The

initial slope of the stress-strain curve in crossply and cloth- reinforced glass composites), Fig.2, shows the expected de- pendence on fibre orientation. Values range from about

107 lb in-2 for the uniaxial longitudinal lay-ups to 5-6 x

106 lb in 2 for the 0°/90° crossply to 2-5 x 106 lb in"2 for the woven cloth composites. The polybenzimidazolematrix composites developed much higher moduli than any

of the other cloth-reinforced materials at cryogenic temperatures. Also, the glass-cloth phenolic composites are found

to have, on the average, slightly higher moduli than glass cloth-epoxies, while glass cloth-polyesters appear as good as the epoxies. The silicone, polyurethane, and Teflon-matrix composites displayed the lowest moduli with an indication

that phenyl silane-matrix composites are of intermediate modulus.

Again taking 77 K as a criterion temperature, the uniaxial

glass-epoxies showed a modulus range of about 8—1 1 x

106lb in"2 with the higher values reflecting variants of S-HTS/Epon 828 66 and Hi-Stren/Resin 2. 4 The 0°/90°

crossply data ranged from 3—7 x 106lb in"2 with values of 5.5-7 x 106 lb in"2 being reported for S-901 glass with a

series of epoxy resin formulations. 3 The cloth-reinforced

330

25

CRYOGENICS . JUNE 1975
epoxies yielded moduli from 2—5 x 106 lb in"2 with the comparatively highest values being reported for 181 glass/ Epon 828 formulations. 2

Although the average moduli of the glass cloth-phenolic

composites was higher than that of the epoxies, literature values ranged from 3—4.7 x 106 lb in"2 which suggests

that no significant difference in moduli should be expected

for good composites made with either epoxy or phenolic

matrices. A detailed look at the glass-cloth polyester data, however, shows a relatively narrow modulus range of about 3.5-4 x 106

lb in"2 suggesting that the strongest composites

made with this polymeric matrix are probably weaker in modulus than the strongest epoxies or phenolics by

about 106

lb in" 2. In a similar way, the glass cloth-silicone matrix composites are still weaker having moduli which vary from 2.5-2.9 x 106 lb in*2 at 77 K. Somewhat fewer moduli

data are available for polybenzimidazole, Teflon, polyurethane, and phenyl silane-matrix composites, but that which

is available suggests that the Teflon matrix produces moduli

of only 1.6-3 x 106lb in"2 with 3.3 x 106 lb in"2 and

3.6 x 106

lb in"2 for the polyurethane and phenyl silane matrices, respectively. The polybenzimidazole data were

obtained with a 181 /polybenzimidazole composite 2 and

are noteworthy not only for the high average value of 4.38 x 106 lb in"2 developed at 77 K, but also for the indication of a substantial increase to 4.9 x 106 lb in"2

at 20 K. As with the ultimate tensile strength, Fig.2 indicates that the moduli become erratic below 77 K. Reference to Fig.5b shows that, with the exception of a few crossply data, the

literature indicates that cooling from 295 K to 77 K will produce a modulus increase of about 0.3-1.2 x 106

lb in"2 with about 0.7-0.8 x 106 lb in"2 being most likely. There

does not appear to be a strong dependence on lay-up. Results of further cooling to 20 K are more difficult to

interpret. In general, the data seem to cluster around a small increase in modulus up to 0.6 x 106

lb in"2 for cloth reinforcement and suggests that a somewhat larger increase on the order of 0.6-1.3 x 106 lb in"2 could be expected for crossply and uniaxial composites. Nevertheless, the scatter from -0.8 x 106 lb in"2 to +3.2 x 106

lb in"2 modulus

change is indicative of something erratic occurring below

77 K.

Again, the only surprises in the 200 K data of Fig.2 are the high value of the modulus of the flexible polyurethane

Adiprene L-100 compared to the room temperature modulus and the almost equally large increase in the polybenzimidazole data.

Composite flexure strength and modulus

Flexural tests are frequently used for screening a large number of composites during development studies, as such

tests are simple and relatively inexpensive compared to tensile testing. In this application, flexure tests have the added

advantage of testing the matrix as well as the reinforcement

fibre. Unfortunately, the state of stress is continuously

changing throughout the flexure specimen as the test proceeds, which makes engineering interpretation of the data

difficult. Consequently, data on flexural strength and

moduli are generally considered valid only for establishing

relative performance ranking.

Fig.3 shows the flexural strengths of the glass-epoxies to be

higher than that observed in tension by approximately

Uniaxial

i 1 -

- (Ref 2,3,31,32,34,3.5,4,5,5.1,66)

0°/90° . -

- (Ref 3,31^.2,34^.5,4,5,51)

90

807060

50 40

H30

(Ref 1,1.2,1 3,2,2 13,3 132,3435,

28,34,35,54,55,57,66,80,97)

Minimal data

_1_

E 30 °>0

50 100 60 200

Temperature, K

250

- 25

20

- 10

- 5 300

Fig.2 Initial tensile modulus of glass-reinforced composites

1 — epoxy; 2 — polyurethane; 3 — phenolic; 4 — polyimide; 5 — polyester; 6 — silicone; 7 — phenyl silane; 8 — Teflon; 9 — polybenzimidazole

400

300

200

1 00

200

150

100

.200

(Ref 2,3,3.1,3.4,35,4) NOL Ring/

0°/90°

(Ref 3,3.1,34,3.5,4,66)

(Ref 1,1.1,1 2,1 3,23,3.1,3 2,3.4,3 5, 121,71,80,89,93,97; Cloth

150 200

Temperature, K 300

Fig.3 Ultimate flexural strength of glass-reinforced composites

1 — epoxy; 2 — polyurethane; 3 — phenolic; 4 — polykmide; 5 — polyester; 6 — silicone; 7 — phenyl silane; 8 — Teflon; 9 — polybenzimidazole

CRYOGENICS . JUNE 1975

26

331
100 k lb in*2 k lb in"2 in the uniaxial specimens and by about 50 in the 0°/90° or cloth lay-ups (verified by comparison of a

1" and tHTM data from the same authors testing the same

composites.) As in the tensile results, the data show the epoxy-matrix

composites to be superior in flexural strength at all temperatures, although the polybenzimidazole-matrix composites are almost as good. Among the other matrix types, the agreement with the tensile data is less clear. For those composites for which there are a reasonable number of data, for examples, the polyester, phenolic, and Teflon-matrix

types, the strength order is the same at 77 K as it is in tension: however, the relative strength difference bear little relationship to the tensile data. The polymide and phenyl

silane-matrix composites rank near the top in flexural strength, while appearing near the bottom in tension. Conversely, the polyurethane-matrix composite appears good in flexure but poor in tension. However, as the data on the former two composites are based on only one or two refer- ences and on a comparison between composites of different authors, caution is necessary in interpreting the results. The

polyurethane-matrix data does reflect the same composite

tested both in tension and in flexure. 2

Within the uniaxial data, a separation has been made between

data generated from flat flexural specimens and those obtained from curved segments of NOL rings because flexural properties obtained from each type of specimen are distinctly

different. Examining the available flexural strength data in more detail, we find that the flat-specimen uniaxial strength of the

glass-epoxies ranged from 325^70 k lb in"2 at 77 K with

the highest values reported for S-901/E-787. 3 A value of 375 klb in"2 was reported for S-901 /Resin 2

4 The NOL speci- men data were significantly lower, ranging from 200—270

lb in"2 at 77 K.4 Flexural strengths varied among the 0°/90° epoxy data from 145-260 k lb in"2 at 77 K with

the highest values again reported in S-90 1 glass using either E-787 or an experimental epoxy formulation. 3 Data for the cloth-reinforced epoxies showed a spread of 95— 175

klbin"2

. The highest values were obtained with 1581/E-787,

with almost as high values reported for 181 glass in a varia- tion of Epon 826 resin. 2 These were the same composites

for which high tensile values were reported. The glass-phenolic composites ranged from 70— 1 10 k lb in"2

at 77 K. The highest values were reported for 181/CTL-91-

LD. 1 Glass-polybenzimidazole data relect only the average

data with 181 glass reinforcement. 2 The glass-polyesters showed a slightly higher range, 80—127 k lb in"2 with the

highest value reported for 181 /Hetron 3 1 ? Reported

flexural strengths of cloth-reinforced Teflon-matrix composites varied from 30-70 k lb in"2

, the highest values being developed with 181/FEP. 2

Following a pattern which is found to repeat itself in all strength properties of glass-polymeric composites, the

flexural strengths all initially increased upon cooling from

295 K to 77 K: however, they then changed in erratic ways upon additional cooling to 20 K. Changes in flexural strength during cooling as reported in the literature are summarized on Fig.5c, which shows that the expected

strength increase from 295 K to 77 K is about 50-80

k lb in"2 for crossply and woven-cloth lay-ups. However,

strength increases of up to 250 k lb in"2 have been reported

for uniaxial composites, suggesting that the magnitude of

the increase is lay-up dependent in flexural strength testing.

A comparison of Fig.5c with that of 5a shows a much

greater scatter in the flexural data as compared to the

tensile data on cooling to 77 K. Upon cooling further to 20 K, Fig.5c indicates that one may

obtain strength changes varying from -50 k lb in"2 to +150 k lb in"2 with a higher probability of a decrease than

an increase. The data do not appear to be lay-up sensitive

at 20 K. The flexural modulus data, Fig.4, show a value of about

8.5 x 10 6lb in-2 for uniaxial glass-epoxy. This is lower

than the average moduli in tension; however, these data

reflect only data for a flat specimen of S-901 /E-787. 3 A check of the data shows that this specific composite had an initial tensile modulus of only about 8.8 x 106

lb in"2

, which suggests that the two methods are giving about

the same answers for the uniaxial case. The same is true

for the 0°/90° case, about 5 x 106

lb in"2 being obtained in both the flexural and tensile modes of testing. The cloth- reinforced polymer flexural modulus data range of 2 3 x

106 lb in"2 to about 5 x 106lb in"2

is also similar to that of the tensile modulus.

Except for the Teflon-matrix showing the lowest modulus,

the silicone-matrix showing next lowest, and the polybenzi- midazole matrix showing one of the highest, in both tests there appears to be little correlation between the relative modulus ranking in flexure and in tension for the same

series of composite types. This is not a condemnation of

the flexural test - it may equally well indicate that average data from the literature cannot be used to predict tensile behaviour from flexural data with any degree of reliability. Examining the flexural modulus data in more detail, we

find that the reported 0°/90° crossply glass-epoxy data

(Ref 3) 50

IOO 150 200

Temperature, K 300 0

Fig.4 Initial tensile modulus of glass-reinforced composites

1 — epoxy; 2 — polyurethane; 3 — phenolic; 4 — polyimide; 5 — polyester; 6 — silicone; 7 — phenyl silane; 8 — Teflon; 9 — polybenzimidazole

332

27

CRYOGENICS . JUNE 1975
ranged from 4.6-5.8 x 106 lb in"2 at 77 K with the highest

values appearing for S-901/E-787. 3 Among the cloth- reinforced composites, the epoxy-matrix data varied from

2.6—5 x 106

lb in'2 with maximum values reported in 181/Epon 828 1 and in 181 glass with modified Epon 828

resin. 2

The phenolic-matrix data encompassed 1.2-5.3 x 106

lb in"2

at 77 K, the highest values being reported for in Conolon 516. 41 -97 The silicone-matrix type composites

showed a relatively small spread of 2.6-3.2 x 10 6 lb in"2

at 77 K, the highest value being reported for 181 glass/ Trevarno F-131. 2 Among the composites for which fewer

data were available, a comparatively high flexural modulus

was reported by Chamberlain, et al 2 for the phenyl

silane-matrix composite Narmco 534. These data again show

high moduli at all temperatures for 181/polybenzimidazole

composites. 2 The polyurethane and polyimide-matrix

flexural modulus data reflect only one reference each. 89 - 2

Again, cooling below 77 K causes the data to become erratic. Fig.5d summarizes the glass-epoxy flexural modulus data and shows that, while on the average one would expect an in- crease of about 0.4 x 106 lb in"2 on cooling to 77 K and a

like increase in further cooling to 20 K, one may find changes

ranging from —0.8 to +2 x 106lb in"2. The data suggest

that the crossply might be more adversely affected by cool- ing to 77 K than the cloth-reinforced specimens; however,

there are not sufficient data to verify this indication.

§j Uniaxial

SJ 07 90°

Cloth

Act 295 K-77 K 77 K-20K

3 H*l . 12 xlO lb ID

A£'

295 K—77 K

-4 0 , xl0

4

lbin"'

A£f 77K-20K

-0.8 0 0.8 1.6 24 -0.8 0 0.8 1.6 24 3.2

x I06

lb in"2

Aa

295 K-77 K

xlO6lb in" 2

kkl M4*kl ...HHH.B.W

0 4 8 12 16 20 24 C xl04 lbin"

2

A£"

295 K —77 K

77 K-20 K

Fl H F

kl+Ut+1 \M , m SI , SI -4 0 4 8

xl04lbin"

2

A£ T 77 K-20K

n , , n m n , n , -0.8 0 0.8 1.6 2.0

(j x I06

lb in" 2 -0.8 0 0.8 1.6 xlO6

lb in"2

Fig.5 Histogram illustrating changes in tensile and flexural properties of glass-epoxy composites upon cooling

Composite compressive strength and modulus

The compressive strength data discussed herein were obtained by compressing in the fibre direction or, in the case of cloth-reinforced materials, in the plane of the cloth. As

most composites are used in fairly thin sheet form, the major

problem is one of avoiding failure by column buckling during

the test. Problems are further accentuated in the uniaxial longitudinal case, where slight misorientation of the fibres can substantially reduce the compressive strength. One observes from Fig. 6 that the average of the data report- ed in the literature for the compressive strength of uniaxial glass-epoxy composites is less than half of the ultimate tensile strength average for the same type composites. Yet,

the 0°/90° compressive strength is but slightly lower than

its tensile strength, while the cloth-reinforced data span about the same range in compression and in tension. Consideration should thus be given to the possibility that the

uniaxial data, and to a lesser extent the 0°/90° data, are lower than the true value due to testing problems.

Among the cloth-reinforced polymers, the epoxy-matrix

composites continue to show superiority over all others. Once again, as in the tensile case, the polyurethane-matrix

materials indicate a remarkable transformation from an

extremely low strength at 295 K to one of the strongest of

the group at 77 K. The glass-phenolics continue their reason- ably good performance previously noted in the tensile

results. The glass-polybenzimidazole appears to rank about

average in compressive strength, similar to its performance

in tension. The glass-phenyl silanes appear to rank somewhat

better in compression than in tension, although few data are

available. The glass-polyesters appear to have relatively low

compressive strengths, although they were ranked among the top in tension. Finally, the Teflon and silicone-matrix composites display consistently the lowest compressive

strengths of all materials surveyed.

Considering the uniaxial compressive strength data in more

detail, we found the reported data at 77 K to vary widely from 100—240 k lb in"2 with about equal scatter at the other temperatures. This is almost twice the percentage

variation found in the uniaxial tensile data even though the

latter data were much more extensive. This large scatter very likely reflects the aforementioned problems inherent

in compression testing. The highest value reported at 77 K

for uniaxial compression was 237klbin"2 in S-901/E-787. 3 The 0°/90° data showed much less scatter, ranging from 106—130 k lb in-2 at 77 K, the highest value reported in biaxially filament-wound S-901 glass in DER 332 epoxy. 3 Among the glass-reinforced composites, the epoxy-matrix

data varied from 90-138 k lb in"2 at 77 K with the highest

value reported for a modified 181/Epon 828 composite. 2 Glass-phenolic properties covered a range of 60— 100 k lb in"2

at 77 K, distinctly lower than those of the epoxies. The

highest value was reported for 181/Narmco 506. 1 The

polyester-matrix composites were by comparison, still another notch down in strength with a 35—68 k lb in"2 showing

at 77 K, the highest compressive strengths being reported

for 181 glass with Hetron 31 or Narmco 527 resin. 2 The

19—40 k lb in"2 range of the glass-sihcones and the 25-40

k lb in"2 range of the glass-Teflons at 77 K leaves little doubt

of the weakness of the latter composite types in compression. The data for polyurethane, polybenzimidazole, and

phenyl silane-matrix composites reflect the work of Chamberlain et al. 2

CRYOGENICS . JUNE 1975

2 8

333
As with tensile and flexural properties, cooling below 77 K

produces somewhat erratic results in compression. The

200 K data appear to be in line except for the glass-Teflon

composite for which cooling to 200 K had no apparent

effect. No uniaxial or 0°/90° data were available at cryogenic

temperatures for compressive moduli. In theory, compressive modulus of any perfectly elastic body should equal

the tensile modulus. Thus, we find that the average values among the cloth-reinforced polymers which appear in Fig. 7 bear a striking resemblance to the tensile modulus

data. Again, the phenolic-matrix composites had highest

values with the epoxy-matrix next highest. Teflon-matrices were clearly the weakest with the polyurethanes starting out equally low in modulus but again rapidly increasing its value to the middle of the group. The silicone-matrices were again on the low modulus side. The only significant change from the tensile modulus ranking is the somewhat

lower value of the polyester and polybenzimidazolematrix composites in compression.

The phenolic composites ranged from 4—8.4 x 10 6 lb in"2

in compression modulus at 77 K, maximum value being

reported for 181/CTL-91-LD. 1 The glass-epoxy data showed

much less scatter at 77 K ranging from 4—5.1 5 x 106

lb in-2. The highest reported value for 181 glass/modified Epon 826

composite. 2 The phenyl silane and polyurethane data

reflects only the values for Narmco 534 and 181/Adiprene

L-100, respectively. 2 Somewhat more compressive moduli

data were available for glass-polyester composites, values ranging from 2.5—4.3 x 10 6 lb in"2 at 77 K with highest

values for 181/Polyester C. 66 The silicones not only aver- aged about like the polyesters but had about the same

spread in 77 K values at 2.3—4.6 x 106 lb in"2

. The data on

Temperature, K

Fig.6 Ultimate compressive strength of glass-reinforced composites

1 — epoxy; 2 — polyurethane; 3 — phenolic; 4 — polyimide; 5 — polyester; 6 — silicone; 7 — phenyl silane; 8 — Teflon; 9 — polybenzimidazole

334

b

Cloth »0

5

- / ^-v 35

4

50

25

20 3

2

15

10

I i Minimal data 5

(Ref l.l 1,12,13,2,80,82,97)

1 1 11 : 50 100 150 200 250 300 0

Temperature, K

Fig. 7 Compressive modulus of glass-reinforced composites

1 — epoxy; 2 — polyurethane; 3 — phenolic; 4 — polyimide; 5 — polyester; 6 — silicone; 7 — phenyl silane; 8 — Teflon; 9 — polybenzimidazole

glass-Teflon and glass-polybenzimidazole composites were

provided by Chamberlain et al. 2 In contrast to its tensile modulus behaviour, the 181 glass/polybenzimidazole composite shows a sharp decline in both compressive strength and modulus on cooling from 295 K to 77 K. This is suggestive of experimental errors.

Composite interlaminar shear strength

Interlaminar shear strength is a property unique to composites; it is the resistance to shearing in the plane of the

fibre reinforcement. It is believed to strongly affect

structural integrity, particularly in compression loading. Like the flexural test, the results of interlaminar shear tests evaluate several composite parameters, including resin strength, resin-fibre bond strength, filament distribution, and matrix porosity. As such, interlaminar shear joins the

flexure test in being a valuable indication of overall composite quality but providing little data directly useful in engineering calculations. Its most valuable application may

well be in quality control during composite manufacture.

Nevertheless, it is of interest to consider the published

values for interlaminar shear, if for no other reason than

to compare the results obtained with the different methods.

Interlaminar shear is usually measured either by the guillo- tine method in which interlaminar shear is forced by the imposition of opposing but offset cuts in a flat tensile specimen, or by a dimensional modification of the flat flexural specimen so as to force failure by shear on the

central layers of the composite (short-beam shear). Of the two methods, the latter is most widely used. Unfortunately,

the results obtained by the two methods are not comparable, the short-beam test usually yielding values higher than

that of the guillotine. The overall situation is made still more complex by the understandable desire of some in- vestigators to obtain interlaminar shear data from filament- wound composites prepared by the NOL ring method in which case the flexural method must be used with a short

section of the ring. Because the specimen is not flat, the

CRYOGENICS. JUNE 1975

29
results are not comparable to either of the above methods.

It is for this reason that the interlaminar data appearing on

Fig.8 have been separated according to the various test methods. The largest discrepancy is observed for the uniaxial composites, the values obtained by the NOL short-beam method

being much higher than those in conventional short-beam or

guillotine methods. It is obvious from Fig.8 that one must

be very cautious in comparing interlaminar shear values published in the literature. In interpreting these kind of

data, it is also necessary to take into consideration that the very high values developed in the NOL short-beam test may

also reflect the generally lower void content in filament- wound composites as compared to vacuum-bagged or autoclave-cured flat lay-ups.

In the case of crossply or cloth laminates, only the conventional short-beam flexure test or the guillotine test may be

used. Fig.8 shows that the results of these two test methods

are in reasonably close agreement for the 0°/90° crossply

lay-ups, while the cloth-reinforced composites show lower

values for the guillotine as compared to the short-beam

test mode. The same is true for the uniaxial lay-ups tested by these two methods. The variety of test methods used and the variety of differ- ent epoxy matrices which have been evaluated make it al- most an exercise in futility to attempt to identify those composite types which have the highest values of inter- laminar shear. The S-901/E-787 composite reported by

Toth et al 3 at 14.7 k lb in-2 (77 K) in short-beam flexure has shear strength almost as high as those reported for the NOL short-beam test. The same composite also developed

high values of interlaminar shear when tested by the same

author in a 0°/90° biaxially-filament-wound lay-up. Other

composites which appeared to have relatively high values of interlaminar shear among their group were S-901 glass

in an experimental resin Epon 826/Empol 1040/Z-6077/

DSA/BDMA 5 in NOL short-beam and a 0°/90° lay-up of

Hi-Stren glass in Epon 828/LP-3/Cure agent D resin in conventional short-beam shear.4 The data indicate that Resin 2 provides somewhat lower interlaminar shear

strength than does E-787; for example, Soffer 5 reports a value of 9.3 k lb in 2 for S-901 /Resin 2 in short-beam shear of flat specimens at 77 K and 14.1 k lb in"2 for the same

composite in NOL short-beam shear. Ignoring the NOL flexure data, one observes on Fig.8 that the interlaminar shear values are the highest for the cloth- reinforced composites and lowest for the 0°/90° crossplies with the uniaxial lay-ups in between. The relatively high values for the cloth composites probably reflect the added shear resistance provided by the convolutions in the woven glass cloth. By similar reasoning, the uniaxial composites may develop higher interlaminar strengths than the 0°/90 crossplies simply because the former provides no

distinct lamella along which shear can propagate.

Again, in a repetition bordering on monotonous, the inter- laminar shear strength is found to become erratic upon cooling from 77 K to 20 K, while the 200 K properties appear

to be as expected.

Composite ultimate tensile strain

In view of the high strength of the glass-polymeric composites coupled with their relatively low modulus and absence

of significant plastic flow at rupture, it might be expected

20 \ (Ref- 3,3.1,3 2 ,34,3 5,4,4 1,5) \ - Uniaxial \

_ 0°/90°

\.

- Cloth

NOL _2^x Flexure \

X / Flexure— S,

Guillotine XX \\

X \

-

X \ x V

•* Flexure

Gui 1 1 ot i n e

Flexure"^ —*t>^\

Guillotine

1

1 1 1 • 1

-7 « 7 b

0 50 100 150 200 250 300

Temperature, K

Fig.8 Interlaminar shear strength of glass-epoxy composites

g

(Ref 3,34,4,5,5.1,35)

50 100 150 200 250 300

Temperature, K

Fig.9 Ultimate tensile strain of glass-epoxy composites

that the strains accommodated at fracture would be some- what larger than those of the higher-modulus advanced

composites. This is substantiated by Fig.9 on which it is seen that ultimate tensile strains are in the 1 0"2 range for

glass reinforcement while, as we shall see in the second part of this review, similar data for the advanced composites are

in the 10"6 range.

This relatively high fracture strain is, of course, a direct

reflection of the high fracture strain of the glass reinforce- ment. Tnis has some interesting consequences because, while the fracture strain of some polymeric matrix materials may approach that of glass at 295 K, the ductility of almost

CRYOGENICS. JUNE 1975 30

335
all polymers tails drastically upon cooling to cryogenic temperatures with the result that at low temperatures, frac- turing the matrix occurs well before fracture of the glass

fibres. 6 A dramatic example of this phenomenon is found

during proof testing of filament-wound pressure vessels where the cracking of the polymeric matrix is very audible.

It is surprising that, at least in the case of hydrostatically loaded pressure vessels, such rupturing of the matrix does not decrease the overall rupture strength of the tanks and may even produce an increase. One must emphasize, however that this does not imply that fracturing of the matrix is always acceptable or that it may be tolerated in other

structures.

Fig.9 indicates that at 295 K, the average fracture strain will be similar for unaxial. 0°/90° and woven-cloth glass-epoxy

composites, the strains being on the order of 3.2 x 10 2

. On cooling, however, marked differences develop between

the lay-up types. Ultimate tensile strains in the uniaxial composites (tested in the longitudinal direction) decreased upon cooling to 77 K while cooling the 0°/90° and cloth

lay-ups resulted in increased fracture strain. Detailed exa- mination of the data shows this to be a true behaviour, as

five of the six composites forming the uniaxial data showed

the decline (the exception being S-901/E-787 3

) while all eight of the 0°,/90° and all three of the cloth composite

types showed substantial increases upon cooling to 77 K.

Again, considering 77 K as a reference temperature, the

uniaxial data were found to range from 1.9—5.3 x 10~2

, the

highest being the previously noted work of Toth et al. 3 The

data of Toth is somewhat suspect, as it is the only data

reporting an increase in tensile strain upon cooling.

.Among the 0°/90° crossply data, values showed a relatively narrow spread of 3.2^4.5 x 10~2

, the S-901/E-787 composite

again showing the highest fracture strain. 3 A value of

4.3 x 10~2 was reported for Resin 2. s Cloth-reinforced data ranged from 3.5—5 x 10~2

, the highest ultimate tensile

stain being reported for 1581/E-787. 3

Data on strain to fracture at cryogenic temperatures are rare for other than the epoxy-matrix types. Kerlin et al 64,65 have reported a fracture strain of 2.75 x 10~2 for a Selectron 5003 glass-cloth polyester composite at the same temperature. Also, Toth et al 3 have reported 5.3 x 10"2 for a crossply glass-polyester composite Selectron 5003.

Static fatigue of composites

Under sustained room temperature loading, glass filaments have been found to deteriorate and fracture when subjected

to stresses substantially below that of their normal ultimate

tensile strength. As such, the failure is analogous to stress corrosion in metals. Some published data indicate that glass-filament-wound pressure vessels may undergo similar deterioration," although some indication has also been

obtained that static fatigue of glass-reinforced composites

may pose less of a problem at cryogenic temperature than at 295 K. 6b Static fatigue will be an important factor in any

composite used in superconducting motors and generators

and, in view of the minimal data presently available, fur- ther studies are needed to clarify the magnitude of the problem and to select formulations providing utmost resis- tance to such failure at cryogenic temperatures. T.T. Chiao

of the Lawrence Livermore Laboratories is currently directing an extensive static fatigue programme at room temperature on dead-weight loaded filaments coated with

resin. As this latter method tests a basic structural element

of the composite and permits many specimens to be run concomitantly, it should provide a useful and relatively inexpensive method of evaluating static fatigue at cryogenic

temperatures.

Composite bearing yield strength

Bearing yield strength is a test designed to determine bear- ing stress as a function of the deformation of a hole through

the composite. The load is applied by a pin inserted into

the hole. The intent of the test is to provide information on

the stress that may be sustained across riveted or bolted

joints without loosening the joint. Bearing yield strength is defined as that stress on the stress-strain curve which corres- ponds to a distance of 4% of the bearing hole diameter when

measured from the intersection of a line tangent to the

stress-strain curve at this point and the zero load axis (see

insert, Fig. 10).

Bearing yield strengths for a series of epoxy-matrix composites have been reported by Toth et al 3 while Chamberlain et al 2 have provided data on silicone, polyurethane, and phenyl silane-matrix composites. These data are summarized on Fig. 10. The glass-epoxy composites developed bearing yield

strengths which increased from about 50 k lb in"2at 295 K to 70-75 k lb in"2 at 77-20 K. The uniaxial and

0°/90° data were very similar, while the cloth-reinforced epoxies were lower at 295 K and at 200 K but increased

their strength to equal that of the others at 77 K. Data on the other matrix types were only available for cloth reinforcements. The phenyl silane composite developed a bearing strength equal to that of the epoxies, while the

Fig. 10 Bearing yield strength of glass-reinforced composites

336 31

CRYOGENICS . JUNE 1975
silicone and polyurethane-matrix materials were distinctly

inferior.

Again, there is indication of erratic behaviour on cooling from 77 K to 20 K.

Cyclic fatigue of composites

Cyclic fatigue of glass-reinforced composites at cryogenic

temperatures has been studied by Brink et al

1 and by

Chamberlain et al 2 in tension and by Fontana et al 34 using a reciprocating beam. Lavengood and Anderson 70 have contributed data on torsional fatigue. The data are not extensive, reflecting the very high cost of generating

complete S-N curves at cryogenic temperatures. Yet, such

testing is mandatory to provide assurance that composite

components will fulfill their life expectance in such applic- ations as rotating cryogenic machinery. Every effort must

therefore be made to restrict such testing to composites

which have the best chance of developing superior fatigue

properties. The data reviewed herein provide some sense of direction; however, there remains an urgent need for some type of relatively inexpensive screening test which

will permit relative ranking of composite cyclic fatigue

performance at cryogenic temperatures.

In comparison to that of the advanced composites or to many conventional alloys, the dynamic fatigue properties of glass-reinforced composites are relatively poor. This is primarily due to the fatigue behaviour of glass-polymeric

materials being controlled by the properties of the matrix. Even at room temperature, the strain accommodation of

the most epoxy resins is less than that of the glass reinforce- ment, and as it is the latter which controls the ultimate

strength of the composite, cracking of the matrix will occur

at ultimate loads far below the composite ultimate strength, allowing corrodents to attack the glass. In crossply lay-ups, such cracking may occur at stress levels as low as 20% of the ultimate composite strength. Lavengood has pointed out that since fatigue life of a glass- reinforced composite is related to the strain capability of

the matrix, embrittling of the matrix due to lowering of the temperature should lower the fatigue life. However, this is not always experimentally verified. Based on an analysis of experimental torsional fatigue data, Lavengood concluded

that fatigue life at cryogenic temperatures is strongly affect- ed by the composite interfacial stress which arises due to

differential thermal contractions of the filament and the matrix. Cooling a composite increases the compressive

forces at the interface and improves fatigue life. Where

fatigue strains serve to increase the compressive interfacial

stresses, the fatigue life may further increase, the reverse being true when strains decrease the interfacial stress. The published tensile-fatigue data of Brink 1 and

Chamberlain 2 and their co-workers were, with one exception, obtained on 181 cloth-reinforced composites. As an

initial criteria, composites were screened by testing at 30%

of their ultimate tensile strength developed at 295 K and

77 K. Those composites failing to achieve 106 cycles were dropped while those which were successful were further

tested at 200 K and 20 K. The conventional polyester- matrix composites 181/fJetron 92, 181/Hetron H-31, 181/

paraplex P—43, and 181/Narmco 527 were unsuccessful as was the silicone-matrix composite 181/Trevarno F— 131.

It is beyond the scope of the present task to comprehensively review these data for which the reader may refer to the

original references. However, in order to provided the reader with a sense of the fatigue performance of glass-polymeric

composites, we have prepared Fig.l 1 from the data of

Brink 1 and Chamberlain. 2 This figure presents the fatigue

strength of the various successful materials as a function of temperature after 106 fatigue cycles, the maximum studied

in this work. It is also instructive to know the percent of

the relevant ultimate strength retained by each composite

type at each temperature after 106 cycles; consequently,

these data also appear on Fig. 1 1

.

Fig.l 1 shows that the absolute magnitude of the stress re- quired to induce fatigue failure at 10 6 cycles generally in- creases with decreasing temperature for all materials studied, only the high-temperature polyesters (Laminae 4232 and

Vibrin 135) and perhaps the silicones (Narmco 513 and

Trevarno F-130) not showing an appreciable increase below

200 K. At all temperatures, these data show that the poly- benzimidazole composite demonstrated comparatively superior fatigue performance both in terms of absolute retained

tensile strength after 106 cycles and in the percentage of the

original strength retained. The epoxies also looked comparatively good, but primarily because of their somewhat higher

initial strength — their percentage retention of strength was among the lowest of the group. The highest values were

reported for Epon 828/DDS.2 The polyurethane, phenyl

silane, and phenolic-matrix composites appeared to group

into an intermediate performance class while the polyester,

silicone, and Teflon-matrix composites showed distinctly lower fatigue properties. Data at 77 K only were available

for a Scotchply 1002 uniaxial glass-epoxy composite, 2 which Fig.l 1 shows to have a much higher strength after 106 fatigue cycles at 77 K than any of the cloth-reinforced

types; nevertheless, even this composite showed retention of only about 52% of its original strength at that temperature. Does the data indicate the existence of a stress limit below

which fatigue life is essentially infinite? In most cases, test- ing was not carried out to a sufficiently large number of

50

30

20

9". %o-'u at temperature 295K 200 K 77 K 20K

Epoxies -I 30 32 37 37

Phenolics-2 41 Polyesters -3 33

Silicones-4 35

Phenyl silanes-5 32

Polyurethanes-6 60

FEP teflon-7 50

«PBI-8 48

Scotchply-9*

42

354245

64

44

36

315740

355252

4140

66

36

2857

+ High temperature types

1 Uniaxial all others 181 glass reinforced

45

40

35

- 30

Polybenzimidozole

I I L

25

20-

10

50 100 150 200 250 300

Temperature, K

Fig.l 1 Cyclic tensile fatigue strength of glass-reinforced composites

after 10 cycles

CRYOGENICS . JUNE 1975

32

337
cycles to answer this question. However, judging from the shapes of the S-.Y curves developed with 181 Epon 1001 by Brink. 1 it appears that this glass-epoxy formulation may

reach such a limit at about 30 k lb in-2 for temperatures be- low 77 K (~ 30% UTS). Conversely, the S-.Y curve for the

181, Epon 828 composite showed no indication of flattening out at 10 6 cycles. With the possible exception of the

silicone-resin composites. 181 Narmco 513 and 181 Trevarno

F-130, there was no evidence that cooling the composites

had any effect on establishment of a fatigue limit where none was evident in the room temperature data. Again, one notices an apparent tendency for the data to become erratic when cooling from 77 K to 20 K.

Impact strength and fracture toughness of composites

Impact strength and its more sophisticated partner, fracture toughness, are measures of the amount of energy which can

be stored in a structure before the energy is released by

fracture. As it is only in recent years that fracture mechanics has been put on solid theoretical grounds and exploited

in homogeneous metals, it is not strange to find that applic- ation of parallel concepts to composites is still in its infancy.

Furthermore, there is a major problem in applying basic

principles of fracture mechanics to composites where, in most cases, multiple cracking occurs so that a unique 'crack

length' cannot be defined. Measurement of the energy required to fracture specimens,

the impact test, is the most simple method of obtaining data on relative material toughness. Such tests show the

glass-reinforced composites to have much greater toughness

than do the advanced composites, probably reflecting the much larger strain-accommodation of glass filaments in

Temperutire, K

Fig.12 Thermal contraction of glass-fibre reinforced composites

338

comparison to the advanced fibre reinforcements. Unfortunately, the literature contains few low temperature data even on this simple parameter and those which are available

are impossible to systematize due to differing test methods,

specimen design, and filament orientation. Uewis and Bush 4 have evaluated a series of epoxies and modified epoxies

reinforced with Hi-Stren glass using the Izod impact method.

For unidirectional composites, they find the 295 K impact

strength to vary from 82- 1 28 ft lb" 1 of notch. On cooling

to 77 K, the measured values ranged from 67- 1 62 ft lb in"1

, some formulations showing 45% increase in impact strength, while others showed as much as a 26% decline. A 0°/90°

crossply test with the same composite formulations yielded room temperature impact strengths from 49—96 ft lb in" 1 and 77 K values of 59—76 ft lb in" 1 with strength changes

ranging from +36% to —26%, no consistent behaviour

being shown by any specific composite formulation. Data

published by other authors, 71,80,103,109,110 ^ equally

confusing. Perhaps part of the answer is provided by

Levin 71 who followed the change in impact strength of

glass-phenolic and glass-epoxy specimens at closely spaced

intervals from 295 K to 77 K and found that the impact

properties peaked sharply at about —30°C followed by a rapid decline at lower temperatures. If the impact properties are indeed such a rapidly changing function of temperature, it might account for the lack of systematic change noted in data taken only at tw7 o or three temperatures, early, this is an area which must be given more attention

in the future if composite behaviour at cryogenic temperatures is to be understood well.

Thermal properties

Thermal contraction of composites

All of the glass-reinforced polymers contract when cooled. As we wish to review the behaviour of such composites on

cooling from 295 K. the data have been plotted as thermal

contraction, avoiding negative values of AL/L. Thermal contraction is very dependent on the type of composite lay-up as well as on the orientation of the composite.

Figs 12 and 13, therefore, present data on values obtained normal and parallel to the fibre reinforcement for uniaxial, 0°/90° crossply and cloth lay-ups. Other factors affecting thermal contraction are the specific resin used and the composite fibre density.

The thermal contraction is always found to be higher in a

direction normal to the fibres than in the fibre plane, the

difference increasing as the temperature is lowered. This

reflects the much higher thermal contraction of the polymer

as compared to the glass reinforcement. This is very evident

in the data for uniaxial epoxy-matrix composites, where at 20 K the thermal contraction normal to the fibres varies from about 2.5 to 5 times that in the fibre direction at 20 K. An apparent anomaly appears in the 0°/90° crossply data, which shows a remarkably small difference in contraction between these two directions. This occurs because, in a filament-wound crossply composite, each layer of glass is in close proximity to the orthogonal layer next to it and

thus provides resistance to dimensional change in the thick- ness direction. This effect is much less evident in the cloth

lay-ups because the latter are prepared by methods which

produce a much lower filament density than is obtained in filament winding.

CRYOGENICS . JUNE 1975

33
30Q »^

100 150 200

Temperature, K

250 300

Fig.13 Thermal contraction of glass-cloth reinforced composites

1 - epoxy 85 - 101

; 2 - polyurethane85 ;3 - phenolic 14

; 6 - phenyl

silane 14

; 7a - Teflon TFE 14

; 7b - Teflon FEP80

; 8 — polybenzimidazole 16

As the thermal contraction of the composite is dominated

by the properties of the polymeric matrix on cooling, one

might anticipate considerable variation in contraction even

within a given matrix type, reflecting, for example, the

effect of additives to epoxy resins. This is found to be true; for example, the AL/L data for epoxy matrix composites

varied from 20-75 x 10"4

in the thickness direction and from 4.4-1 1.5 x 1(H in the fibre direction. In the 0°/90°

lay-ups, the epoxy composites showed variations of 5.1-13.8 x 10A in thickness and 9.2-13 x 10-* in the

fibre direction. The polyester resins gave similar data spread

in the 0°/90° lay-ups. A much greater variation is observed in the thermal contraction among cloth-reinforced composites, reflecting both the

larger variety of matrix materials for which data are available and the greater variation in fabrication method. The glass- Teflon composites have the highest thermal contractions of

all the materials examined. Furthermore, the Teflon matrix may be either of the TFE or FEP type which have markedly

differing thermal contraction characteristics. This difference does not appreciably affect contraction in the fibre plane due

to the restraint provided by the fibres. However, the differ- ence is marked in the direction normal to the fibres. As

indicated by the dashed curves 7a and 7b on Fig.13, the TFE-matrix composites have twice the transverse thermal

contraction of the FEP-matrix types. The epoxy and the polyurethane-matrix composites also have relatively high thermal contractions. Unfortunately, no data were available in the fibre normal directions for these composites. The data indicate that the polyester-matrix

composites are next in decreasing order of thermal contraction, with the fibre normal data being approximately 2.8 times greater than in the fibre plane. Showing the least ther- mal expansion in the fibre plane are the phenolic, silicone, phenyl silane, and polybenzimidazole-matrix composites.

Of these, the phenolic composite data appear to be abnorr mally high for the fibre normal case, since, if all else were

equal, the order in the fibre normal direction should follow

that in the fibre plane.

All of the polymeric matrices create compressive stresses in the glass reinforcement upon cooling to cryogenic temperatures, such stresses being added to those already existing from cooling from the cure temperature to room temperature. This is primarily manifested as shear stress at the

resin-fibre interface, so that it might be expected that an

effect would be seen on those strength properties which

are sensitive to debonding of the matrix-glass interface. One

such property is fatigue, and from Fig.l 1, it is noted that the fatigue properties of the Teflon-matrix composites are poor, while that for the polybenzimidazole is good, consis- tent with the respectively high and low thermal expansion

of these two composite types. A similar relationship holds

for the flexural strength, Fig.3. Interfacial residual stress is but one of the factors influencing fibre debonding, others being the ability of the interface to sustain a shear load and

the strength of the matrix itself.

Specific heat of composites

The specific heats of glass-reinforced composites show an

almost linear dependence on temperature from 295 K to 77 K, and relatively small difference between matrix types. The specific heats are relatively high compared to most

metals, being roughly comparable to that of aluminium but

substantially above that of titanium, iron, or copper. Teflon and polyester-matrix composites have the highest specific heat. The epoxy, polybenzimidazole, silicone, and phenolic- matrix composites are bunched on the lower specific heat

side of the group. The phenyl silane-matrix composite seems

Temperature, K

Fig. 14 Specific heat of glass-reinforced composites

1 — Epoxies 15,16

; 2 — phenolics 15; 3 — polyester 15; 4 — silicones 1 5 — phenyl silanes 15

; 6 — Teflon 15; 7 — polybenzimidazole 16

; 8 — phenyl formaldehydes 72

CRYOGENICS. JUNE 1975

34

339
I I I I I l_l 0 50 100 150 200 250 300

Temperature, K

Fig. 15 Thermal conductivity of glass-fibre reinforced composites

to have a slightly larger temperature dependence, starting out in the middle of the group at 295 K, but showing the lowest specific heat at 77 K. The data reflect the work of Campbell et al 15,16 with the

sole exception of the phenyl formaldehyde contribution of

Luikov. 72 Except for the epoxies, the data reflect only one

composite for each type of matrix. The epoxy data reflect the average of three compositions: a unidirectional and a crossply YM-31-A/DER-332 and a unidirectional S-994/ E-787 composite. The spread in specific heat values for these three formulations at 295 K, 200 K, and 77 K was

8.16-9.21, 6.07-6.7, and 2.09-2.72 x 10"2 J g-'K 1

, respectively, with the highest value associated with the S-994/E-787 composite.

Thermal conductivity of composites Thermal conductivity is also an important parameter in cryogenic design. Unfortunately, improper experimental

technique invalidates some of the data in the literature; primarily, failure to properly compensate for radiation

losses, which can introduce errors approaching 100-200%

in the higher temperature ranges. Figs 15 and 16 presents what the author believes to be valid data. As with thermal contraction, conductivity is dependent on

the type of lay-up and on the orientation of the composite.

Figs 15 and 16 therefore present data on values obtained normal to the fibre reinforcement (thickness direction) and

in the plane of the reinforcement for uniaxial, 0°/90°

crossply and for woven-glass cloth lay-ups. Also, as with

contraction, the thermal conductivities are affected by the type of matrix resin, necessitating a differentiation on this basis in the figures.

Unlike the other thermal properties of composites, the thermal conductivity is affected by the ambient atmosphere

in which the measurement is made, that is it differs in helium, nitrogen, or in vacuum. The literature indicates that, compared to values obtained in helium, data taken in nitro- gen will average 7% lower, while in a vacuum the data will, on the average, be lowered 20%. Campbell et al 16 suggests

that the lower value in vacuum primarily reflects the diffi- culty of obtaining good contact between the composite

specimen and the mating parts of the thermal conductivity apparatus when operating in a vacuum. An ambient atmosphere of nitrogen or helium increases the measured conductivity by reducing the contract resistance; however,

diffusion of the gas (particularly of helium) into pores within the composite also contributes to an increase as the gas provides an overall improvement in the thermal path within

the composite. This explanation is consistent with the

effect of different ambient atmospheres being observed

over the entire 20—295 K temperature range. The present author believes, however, that consideration should also be

given to the effect of matrix cracking at very low temperatures which should act to increase the difference between

values measured in the various ambient media. The data appearing in Figs 15 and 16 are averages of data taken in all three media. The difference between the fibre normal and the in-plane conductivities for a given composite and the absolute spread of conductivity values among the various composite classes

is widest at room temperature, converging as the temperature is lowered, consistent with a theoretical zero value for conductivity at absolute zero temperature. The conductivity at 20 K varies from 30—50% of that at room temperI I I I I L_ 0 50 100 60 200 250 300

Temperature, K

Fig. 16 Thermal conductivity of glass-cloth reinforced composites

1 — epoxies; 2 — phenol ics; 3 — polyesters; 4 — silicones; 5 — phenyl silanes; 6 — Teflons; 7 — polybenzimidazoles

340 35

CRYOCE ICS . JUNE 1975
ature for most of the composites. Exceptions are the cloth- reinforced silicones, curve 4, Fig. 16, in which the 20 K value was only 15% of that at 295 K and the cloth-reinforced Teflon, curve 6, Fig. 16, for which the 20 K value was re- ported to be 70% of the 295 K value, both of these exceptions being for the fibre normal direction.

Fig. 15 contains data for uniaxial and 0°/90° epoxy lay-ups and for a 0°/90° silicone-matrix lay-up. The highest ther- mal conductivity was for the filament-wound uniaxial epoxy composite in the plane of the fibre. This is to be expected from the high density of continuous glass fibres

in that direction. The conductivity in the fibre normal

direction is 25—50% less, reflecting the lower thermal conductivity of the matrix. Thermal conductivity in the fibre plane of the 0°/90° epoxy lay-up is only about 80% of that

for the uniaxial composite, reflecting the lower effective

fibre density in the direction of the heat flow.

Fig. 15 shows the 0°/90° silicone-matrix composite to have

a substantially lower in-plane conductivity than the 0°/90° epoxy lay-up. This does not imply that the conductivity of

the silicone matrix is less than that of the epoxy; indeed,

the marked similarity between the fibre normal data for both

polymeric matrices indicates that the matrix conductivity is very similar in both materials. The difference reflects the

relatively inferior heat transfer properties of glass roving

(fibre bundles) which was used for the silicone-matrix composite as compared to the continuous filament used with

the epoxy. Among the cloth-reinforced composites, the highest thermal

conductivity was found for the silicone-matrix composites.

Only fibre normal data were available; however, even in this least conducting orientation, the conductivity exceeded that of all the other materials in their more favourable in-plane

direction. The polybenzimidazole composite was next

lowest, and in decreasing order were the epoxies, phenolics, and phenyl silanes, all with approximately equal conductivities, both in-plane and plane normal. The lowest conductivities were evidenced by the polyesters and the Teflons. The only composites for which there was more than one

literature reference were the epoxies. For the latter, the data spread was found to be much greater for the in-plane

conductivity than for the plane normal, as would be ex- pected in view of the relatively high conductivity of the

fibres compared to the matrix. As an example, the data which averaged 0.50 W nr 1 K"1 for the 295 K cloth-rein- forced epoxy ranged from 0.30 to 0.65 W nr1 K" 1

, while

the plane normal data for the identical composites average 0.35 W nr1 K"1 but ranged from 0.30 to 0.40 W nr1 K 1

. The scatter decreased significantly at lower temperatures.

Comments on glass-reinforced composites

We have reviewed the properties of filamentary-glass rein- forced composites at cryogenic temperatures in order to provide the reader with an overall feel for their behaviour and the magnitude of the properties which may be expected.

Having this information, for what applications should glass- reinforced composites be used? What composite formulation should be selected?

It would appear that glass-reinforced composites are most

useful in applications requiring high tensile strength combined with high toughness and low thermal conductivity, but where stiffness is not required and where cyclic fatigue

is not a major problem. In such applications, glass-rein- forced composites have given and will continue to give excellent performance at relatively low cost. Filament- winding techniques should be used whenever possible in order to obtain the highest quality composites; in particular, the lowest void content. Resin 2 is unique in that it was developed specifically by

NASA for use in glass-reinforced composites at cryogenic

temperatures. This resin was found by Soffer and Molho 5 to be the overall choice among 41 candidate resins, using the

criteria of composite and fibre tensile strength, composite

tensile modulus, ultimate tensile strain, thermal shock resis- tance, coefficient of linear thermal contraction, and inter- laminar shear strength as well as favourable processing characteristics. For this reason, as well as on the basis of the

successful use of this resin in the cryogenic industry, the NASA 2 formulation must be considered as a prime candi- date for general cryogenic use. The reader is cautioned,

however, that the plasticizers used in this resin cause the

resin to become weak at the elevated temperatures required

for vacuum degassing of a composite assembly. Components

made with this resin must be adequately supported during

this operation.

Early in their investigation of cryogenic resins, Soffer and Molho 5 rejected the Polaris (E-787, 58-68 R) formulation because its glass-transition temperature of over 1 15°C sug- gested brittleness at cryogenic temperatures. However,

available data show that this resin formulation is capable of producing composite mechanical properties equal to those with Resin 2 at crypgenic temperatures, at least for certain filament orientations. This suggests that low temperature

ductility and thermal shock resistance may not always be of primary significance in selecting a resin for cryogenic applications. Nevertheless, the reader is cautioned that inadequate

data exLt on degradation of composite properties during

fatigue at cryogenic temperatures, and the effects of repeated thermal cycling are not well understood. Ductility may well play an important role under the latter circumstances. In the absence of further experimental data, critical composite

components subjected to dynamic loads or to cyclic ther- mal stresses under service conditions should be tested under

simulated operational conditions, regardless of the selected matrix resin. The present review is restricted to those composites for which

cryogenic property data are available in the literature. The

contemporary purchaser of composite components may

find other types of polymeric matrices recommended by

various fabricators for use at cryogenic temperatures. The

present review should not be construed as prejudicial to any

such recommendation. The purchaser should, however, as- certain that any such recommendation is supported by ade- quate experimental or service data at the temperature of

interest. An unexpected result of the present review is the surpris- ingly good overall performance of the glass-polybenzimi- dazole types. Available data indicate that at 77 K, such a composite ranks second only to the epoxies in ultimate tensile strength and in flexural strength, while developing tensile and flexural moduli superior to that of the epoxies. Compressive properties appear only average; however, the data

indicate that the polybenzimidazole-matrix composite with- stands cyclic fatigue stresses with less strength loss than is the

case for the other reported matrix materials, particularly in regard to the percent of original ultimate strength retained

CRYOGENICS . JUNE 1975 36

341
after fatigue cycling. A key may be the comparatively low

thermal contraction of the polybenzimidazole resin as compared to the other polymer types which should reduce re- sidual interfacial stresses between the fibre and the matrix of

the composite.

A pervasive characteristic of glass-polymeric composites is the erratic behaviour of the mechanical properties on cooling from 77—20 K. With few exceptions, the strength properties increase on cooling from 295—77 K; however, on

further cooling to 20 K, the data indicate that such properties may either increase, decrease, or remain unchanged,

the behaviour being quite unpredictable, even among

composites of the same matrix type. It is difficult to attri- bute this simply to the matrix becoming suddenly much

more brittle between 77 K and 20 K, as cooling to 77 K

has already decreased the strain capability of the matrices

to a level far below that of the glass reinforcement. The

present task does not permit more than a cursory consideration of the possible factors involved in this phenomena;

however, the author believes it relevant to call attention to recent studies on the low temperature properties of poly- men which have provided convincing evidence that, in at

least the linear polymers of the polycarbonate (PC) and

polyethelene tereptlialate (PET) types, the media in which

the low temperature test is conducted can strongly affect the fracture mode and the resultant mechanical properties measured for the polymer. 125 ' 126 The experimental

evidence indicates that the failure in such polymers in the cryogenic range is controlled by a crazing phenomena which,

in turn, is related to the activity of the gas or liquid in contact with the polymer surface. Such studies have not been

extended to the strongly cross-linked polymers; however,

until proved otherwise, it must be considered a possibility that the mechanical properties of glass-polymeric composites

may be influenced by the ambient media such that data obtained at 20 K or 77 K in liquid hydrogen or nitrogen

may not be the same as those which would be obtained in

a helium atmosphere.

This review^ suggests that the following work would be of

value in implementing expanded use of glass-reinforced

composites for demanding cryogenic structural applications such as would be encountered in superconducting

machinery:

1. The reason for the erratic mechanical property be- haviour in polymeric-matrix composites below 77 K must

be ascertained. In particular, it is imperative to determine whether or not the type of cryogen in which tests are con- ducted has a significant effect on the mechanical

properties.

2. Material research and evaluation must be directed toward obtaining the ty pe of basic composite cryogenic

property data which will be of most value to the design

engineer. Contemporary composite theory requires full mechanical characterization of a uniaxial lamella of the composite of interest, that is, an experimental determination of the strength amd moduli values are required for the prediction of limiting property values in complex

crossply lay-ups. Accurate tensile and compression data

are required in the longitudinal and transverse modes

plus accurate values for intralaminar shear. Even at room

temperature, tensile and compression testing in the rein- forcement direction has proved difficult to perform with

acceptable accuracy. Testing problems will be further complicated at cryogenic temperatures; however, these problems must be solved.

3. Having come to terms with the problems posed in 1 and 2, the static mechanical property data and the thermal property data for the best of the glass-reinforced composites must be extended down to 4 K. Available data suggest

that the major characterization effort should be made on S

glass in a NASA Resin 2 matrix. S glass should be characterized in the Polaris formulation as well, but with emphasis on its potential special use as uniaxial filament-wound

cryogenic support members. All test materials must be

fully characterized for resin/fibre density, void content, and

fibre alignment before mechanical or thermal investigation.

4. Data on the performance of composites under dynamic

loading conditions at cryogenic temperatures are minimal

to non-existent. Yet, these type of data are mandatory if composites are to be used in cryogenic machinery. The

available data are encouraging in that they suggest that fatigue performance at cryogenic temperatures is generally

superior to that at room temperature. However, this will have to be more fully documented. As high-cycle fatigue

testing at 4 K is very expensive, the materials included in such a testing programme must be carefully selected. For

this purpose, it would be very desirable to have an efficient screening type of test capable of correlating incipient

damage with expected fatigue life. Data on thermal fatigue

are also required, that is, the effect of repeated cool-downs

on both the static and dynamic properties of composites must be determined.

The author wishes to thank Dr R. P. Reed for his consultation and for review of the completed manuscript. The

author also wishes to express appreciation to the NBSN'O AA library staff of the Boulder Laboratories for their assistance in retrieving the many documents required for

this work. This research was supported by the Advanced

Research Projects Agency.

Appendix

The following materials are referred to in this report. [The

trade names used in this paper are those used by the cited authors. Generic names have been substituted whenever

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.]

Glass fibres S-HTS, S-901, 1581,

181, YM-31-A

Hi-Stren Epoxy resins Epon822, 828, 1031,

58-68R

E-.787 ERL 2256

660 FW

DER332

Hienolic resins CTL-91-LD

Owens-Corning Glass Company

Aerojet General Corporation

Shell Chemical Corp, Plastics & Resin Div US Polymeric Corporation

Union Carbide Plastics Company

Stratoglas Div, Air Logistics Corp

Dow Chemical Corporation

Cincinnati Testing Laboratories, Div of Studebaker-Packard Corp

342 37

CRYOGENICS . JUNE 1975
Conolon 506, 507

Phenyl silane resins Narmco 534

Polyester resins Narmco 527

Hetron31,92

Polyester C

Paraplex P43

Laminae 4232

Vibrin 135

Silicone resins Trevarno F-130, F-131 Narmco 513

Narmco Materials Div, Telecomputing Corp

Narmco Materials Div, Telecomputing Corp

Narmco Materials Div, Telecomputing Corp

Durez Plastics Div Hooker

Chemical Corp

US Polymeric Corp

Rohm and Haas Corp

American Cyanimid Corp

Naugatuck Chemical Co, Div of US Rubber Corp

Coast Manufacturing Company

Narmco Materials Div, Telecomputing Corp

Polyurethane resins Adiprene L-100 E. I. DuPont Corporation

Polybenzimidazole resins

Imidite Narmco Division, Whittaker Corp

Flexibilizers, coupling agents, and hardners Empol 1040 Emery Industries, Inc LP-3 Thiokol Chemical Corp

Z-6077 Dow Chemicai Corporation

ZZL-0870 Union Carbide Plastics Company

Miscellaneous materials Cure Agend D polyamine salt DDS diaminodiphenyl sulphone

DSA dodecenyl succinic anhydride

BDMA benzyldimethylamine

NMA nadic methyl anhydride

TFE polytetrafluoroethylene FEP polytetrafluoroethylene copolymer with

hexafluoro propalene

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SpaceiAdministration, Lewis Research Center, Cleveland, Ohio

(1969) (N69-31300) (Identical to SAMPE 15, 249 46 Hanson, M. P., Richards, H. T., Hickel, R. O. 'Preliminary in- vestigation of filament-wound glass-reinforced plastics and liners for cryogenic pressure vessels', NASA TN D-2741, National Aeronautics and Space Administration, Lewis Research Center, Cleveland, Ohio (1965) 47 Hanson, M. P. 'Glass-, boron-, and graphite-filament-wound

resin composites and liners for cryogenic pressure vessels', NASA

TN D-4412, National Aeronautics and Space Administration, Lewis Research Center, Cleveland, Ohio (1968) [Identical to NASA TM X-52350, (1967)] 48 Hanson, M. P. 'Static and dynamic fatigue behavior of glass filament-wound pressure vessels at ambient and cryogenic temperatures', NASA TN D-5807, National Aeronautics and Space

Administration, Lewis Research Center, Cleveland, Ohio (1970) (CFSTI-CSCL-20 K) 49 Haskins, J. F., Hertz, J. 'Thermal conductivity testing of coast F-224-6 phenolic-fiberglass laminate', General Dynamics

Convair Report No AR-592-M82 (1963) 50 Haskins, J. F., Hurlich, A- 'Measured values for the coefficients of linear expansion of Plycel 420 and Conolon 506 at low temperatures', Convair/ Astronautics Report MRG-154 (1960)

51 Haylett, J. W., Rottmayer, E., Butcher, 1. 'Advanced composite material study for millimeter wavelength antennas',

CRYOGENICS . JUNE 1975

40

345
Technical Report AFML-TR-7 1-205, Vol 1, Air Force Materials Laboratory, Wright-Patterson Air Force Base, Ohio (1971) (AD 893 368) 52 Haylett, C. E. 'Advanced composite material study for milli- 71 meter wavelength antennas, Volume II: Environmental tests', AFML-TR-71-205, Vol 2, Air Force Materials Laboratory,

Wright-Patterson Air Force Base, Ohio (1971) (AD 893 358 L) 72 53 Herring, H. W., Baucom, R. M., Pride, R. A. 'Research on

boron filaments and boron reinforced composites', 10th

National Symposium, Society of Aerospace Material and Process Engineers 10 (1966) B-21

54 Hertz, J. 'Tensile testing of Conolon 506 at room and sub- zero temperatures', Convair/ Astronautics Report MRG-120 73 (1959)

55 Hertz, J. 'Tensile testing of Adlock 851, Adlock PG-LA and

Adlock EG-1 1 A-81 A from -423°F to 78°F' Convair/ 74

Astronautics Report MRG 237 (1961) 56 Hertz, J. 'Investigation of potential low temperature insulations', General Dynamics/ Astronautics Report GS/A-ERR-AN-668 75 (1964) 57 Hertz, J. ' The effect of cryogenic temperatures on the mechanical properties of reinforced plastic laminates', General Dynamics Report AR-592-M15 (1963) (AD 405 170) 76 58 Hertz, J. investigation into the high-temperature strength degradation of fiber-reinforced resin composite during ambient

aging'. General Dynamics/Convair Report No GDCA-DBG73-005

(Final Contract NAS 8-27435) (1973) 77

59 Hoggatt, J. T. 'Development of cryogenic PRD^+9-1 filament- wound tanks', NASA CR-1 20835, National Aeronautics and

Space Administration, Lewis Research CEnter, Cleveland, Ohio 78

(1971) (N72-24941)

60 Hoggatt, J. T. 'High performance filament wound composites

for pressure vessel applications', Society of Aerospace Material and Process Engineers, National Technical Conference, Space 79

Shuttle Materials, Vol 3 (1971) 157

6 1 Hust, J. G. 'Thermal conductivity of an epoxy-fiberglass laminate', NBS Cryogenic Division Unpublished Laboratory Note 73-1,(1973) 80 62 Johnston, FL L., Brooks, H. E. 'Impact strength of various metals at temperatures down to 20 Absolute', Ohio State

University Cryogenic Laboratory Report TR 264-17 (1952)

63 Keller, C.W. 'Fiberglass supports for cryogenic tanks', NASA CR-120937 (Final), National Aeronautics and 81 Space Administration, Lewis Research Center, Cleveland, Ohio (1972) (N72-33564)

64 Kerlin, E. E., Smith, E. T. 'Measured effects of the various combinations of nuclear radiation vacuum and

cryotemperatures on engineering materials: Biennial

Report', NASA CR-77772, National Aeronautics and 82

Space Administration, George C. Marshall Space

Flight Center, Huntsville, Alabama (1966) (N66-

35963) 83

65 Kerlin, E. E., Smith, E. T. 'Measured effects of the

various combinations of nuclear radiation, vacuum and

cryotempertures on engineering materials: Annual 84

Report', NASA CR-58830, National Aeronautics and

Space Administration, George C. Marshall Space Flight Centre, Huntsville, Alabama (1964) (N64-33043) 85

66 Keys, R- D., Kiefer, T. F., Schwartzberg, F. R. 'Cryogenics properties of high-strength glass-reinforced 86

plastics', Advances in Cryogenic Engineering 11 (1966) 470

67 Krause, D. R- 'Development of lightweight material composites to insultate cryogenic tanks for 30-day 87

storage in outer space', NASA CR-1 23 79 7, National Aeronautics and Space Administration, George C. Marshall Space Flight Center, Huntsville, Alabama (1972) (N72-30495)

68 Krause, D. R., Fredrickson, G. O., Klevatt, P. L. 'Effects 88

of cyclical environments on high-performance multi-layer

insulation materials', Society of Aerospace Material and

Process Engineers National Technical Conference, Space

Shuttle Materials, Vol 3 (1971) 639 89 69 Lantz, R. B. 'Materials for filament wound cryogenic

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70 Lavengood, R. E., Anderson, R. M. 'Matrix properties 90

controUing4orsional fatigue life of fiber reinforced com- -

posites', Technical Papers, 24th Annual Technical Conference, Society of the Plastics Industry, Section 1 1-E

(1969)

Levin, V. A., Naumenkov, P. G., Shchitov, M. V. 'Some

properties of plastics at low temperatures', Plasticheskia Massy 11 (1966) 64

Luikov, A. V., Vasiliev, L. L, Shashkov, A. G. 'A method for the simultaneous determination of all ther- mal properties of poor heat conductors over the temperature range 80 to 500 K', Proceedings 3rd American

Society of Mechanical Engineers Symposium, Purdue

University (1965) 314

Lyon, D. N., Parrish, W. R. 'Low temperture thermal

conductivities of two high compressive strength

materials', Cryogenics 7 (1967)

Maher, L. E. 'Some problems arising from the use of

hydrogen-fuelled propulsion systems' Journal of the Royal Aeronautical Society 68 (1964) 765 McKannon, E. C, Gause, R. L. 'Effects of nuclear

radiation and cryogenic temperatures on non-metallic

engineering materials', Journal ofSpacecraft 2 (1965) 558

Morris, E. E. 'Glass-fiber-reinforced metallic tanks for cryogenic service', 12th National Symposium, Society of Aerospace Materials and Process Engineers, Vol 12

(1967) Section AS^t (also NASA CR-72224)

Morris, E. E. 'The performance of glass-filament-wound

pressure vessels with metal liners at cryogenic temperatures', Journal ofMaterials 4 (1969) 970

Morris, E. E., Alfring, R- J. 'Cryogenic boron-filament- wound pressure vessels', Composite Materials: Testing and

Design, ASTM STP 460 (American Society for Testing and Materials, 1969) 430

Morris, E. E., Landes, R. E. 'Cryogenics glass-filament- wound tank evaluation', NASA CR-72948 (Final). National Aeronautics and Space Administration, NASA Lewis

Research Center, Cleveland, Ohio, (1971) (N72-14696)

Mowers, R. E., Leib, J. H., Sherman, S. 'Program of

testing nonmetallic materials at cryogenic temperatures', Rocketdyne Corporation Report R-3498, Rocket Pro- pulsion Laboratories, Edwards, California (1962) (AD

294 772)

Nadler, M. A., Yoshino, S. Y.. Darms, F. J. Boron/ epoxy support strut for non-integral cryogenic tankage', North American Rockwell Space Division Report SD

68-99501 (1969) [See also 15th National Symposium

SAMPE (April 1969) and North American Rockwell

Report SD 995 2 (1968)]

Nelson, L. F. 'Compressive strength of Conolon 506 at +75 F and -320 F', Convair/ Astronautics Report No

27E 1336 (1962)

Nelson, L. R, 'Mechanical properties of Adlock 851 at room temperature, 1000 , — 320 and -423 F', Convair/

Astronautics Report No 55E 812 (1961)

Nelson, P. T., Archer, J. S. 'Graphite reinforced plastic EHF antenna', TRW Systems Group, Redondo Beach,

California Report No 99900-7128-RO-ll (1969)

Patten, P. M. 'Internal insulation liner alteration', Douglas Aircraft Company Report No SM 45975 (1964)

Perkins-Elmer Optical Group, Norwalk, Connecticut:

Work-in-Progress on Contract No F33615-72-C-2033,

Air Force Systems Command, Wright-Patterson Air Force

Base, Ohio

Pink, E., Campbell, J. D. "The effect of strain rate and

temperature on the deformation behavior of reinforced and unreinforced epoxy resin', Oxford University Depart- ment of Engineering Report No 1040/72, Oxford, England (1972) (N73-10568)

Pirgon, O., Wostenholm, G. H., Yates, B. "Thermal

expansion at elevated temperatures, IV. Carbon-fibre

composites', Journal ofPhysics D: Applied Physics 6 (1973) 309

Pride, R. A., Stein, B. A., Schmidt, F. W. 'Mechanical

properties of polyimide-resin glass-fiber laminates for various

time, temperature and pressure exposures', Technical Papers, 23rd Annual Reinforced Plastics Technical and Mangement

Conference, Washington, DC Section 17-c (1968) 1 Ratciiffe, E. H. 'Thermal conductivities of plastics with glass, asbestos and cellulosic fiber reinforcements', Applied Material Research 5 (1966) 200

346 CRYOGENICS. JUNE 1975

41
91 Roscland, L. M. 'Materials for cryogenic usage', Technical

Papers, 21st Annual Technical Conference, Society of the

Plastics Industry Section 4-C (1966) 1 92 Roseland, L. M. 'Investigation of structural properties at cryogenic temperatures of filament-wound pressure vessels containing both organic and glass filaments', Douglas Aircraft Corporation Report No SIVM8409, (1966)

93 Ross, J. E. 'Fiberglass laminate - ultimate tensile and flexural strength tests at room temperature, -100°F and 320°F', Convair Astronautics Report No 7E 1687 (1959) (AD 830

230) 94 Sanders, R. H., Weleff, W. 'Final report on GTR-17 effects of radiation on organic materials irradiated in liquid hydrogen',

Aerojet-General Corporation Report No RN-S-0327 (1967)

95 Sanger, M. J., Molho, R., Howard, W. W. 'Exploratory

evaluation of filament-wound composites for tankage of

rocket oxidizers and fuels', AFML-TR-65-381 , Air Force

Materials Laboratory, Wright-Patterson Air Force Base, Ohio

(1966) (AD477 455)

96 Sanger, M. J., Reinhart, T. J. 'Development of filament- wound tankage for rocket oxidizers and fuels', Technical

Papers 12th National Symposium, Society of Aerospace

Material and Process Engineers Section AS-7 (1967) 97 Sewell, J. J., Kuno, J. K. 'Aerospace use of plastic hardware and thermal insulation', Technical Papers, 17th Annual

Technical Papers, 1 7th Annual Technical Conference, Soriety

of the Plastics Industry, Section 7-A (1962) 1 98 Shriver, C B. 'Design and fabrication of an internally insulated filament wound liquid hydrogen propellant tank', NASA

CR-127, National Aeronautics and Space Administration Washington, DC (1964) (N65-10775)

99 Soltysiak, D. J., Toth, J. M. 'Static fatigue of fiber glass pressure vessels from ambient to cryogenic temperatures', Technical Papers, 22nd Annual Technical Conference, Society

of the Plastics Industry, Section 14-E (1967)

100 Speare, J. C. 'Preliminary sizing of filament-wound RNS

tanks', Report No TOR-0066 (5759-07)- 1 3, Space and

Missile Systems Organization, Air Force Systems Command,

Los Angeles Air Force Station, Los Angeles, California ( 1 970) (AD 872 626)

101 Steinhauer, R. A. 'Linear thermal expansion of 828CL 181

cloth laminate', Douglas Aircraft Company Report No MP 1

1

979 (1961) 102 Stinnett, W. D. 'Cryogenic tensile properties of selected

materials', NASA CR-71751, AEC-NASA Space Nuclear

Propulsion Office, Report No 2712 (1964) (N66-22816)

103 Suezawa, Y., Hojo, H., Nakamura, K. 'Impact characteristics of fiberglass reinforced plastics at low temperatures', Kagaku

Kogaku (Chemical Engineering, Japan) 33 (1969) 1051

104 Toth, J. M. 'Barrier films for filament-wound fiberglass cryogenic vessels', Advances in Cryogenic Engineering 1 (1964) 537

105 Toth, J. M., Soltysiak, D. J. 'Investigation of smooth-bonded

metal liners for glass fiber filament-wound pressure vessels', NASA CR-72165 (Final), National Aeronautics and Space

Administration, Lewis Research Center, Cleveland Ohio

(1967) (N67-25070)

106 Toth, J. M., Sherman, W. C, Soltysiak, D. J. 'Investigation of smooth-bonded metal liners for glass fiber filament-wound

pressure vessels', Douglas Missile and Space Systems Division Report No SM^9384, Quarterly Report No 3, Contract No NAS 3 6293, NASA Lewis Research Center, Cleveland Ohio

(1966) 107 Toth, J. M., Sherman, W. C, Soltysiak, D. J. 'Investigation of structural properties of fiber-glass filament-wound pressure

vessels at cryogenic temperatures', NASA CR-54393, National Aeronautics and Space Administration, Lewis Research

Center, Cleveland, Ohio (1965) (N65-35392)

108 Toth, J. M., Barber, J. R. 'Structural properties of glass-fiber filament-wound cryogenic pressure vessels', Advances in Cryogenic Engineering 10 (1965) 134

109 Voloshenko-Klimovitskii, Yu. Ya., Belyaev, Yu. A., L'vof,

B. S., Schpakovskaya, E. I. 'Strength of cold-hardening GRPs based on PN-1 resin under impact tension at normal

(20 C) and low (-196°C) temperatures', Plasticheski Massy 6(1964) 39 110 Voloshenko-Klimovitskii, Yu. Ya., Belyaev, Yu. A., Korenkov,

Yu. A. 'Impact tensile tests on glass-fibre reinforced plastics

at normal and low temperatures', Plasticheski Massy 5 (1963)

51

111 Watson, J. F., Christian, J. L., Hertz, J. 'Selection of materials

for cryogenic applications in missiles and aerospace vehicles', Convair/Astronautics Report No MRG 132-1 (1960)

1 1 2 Weleff, W. 'Effect of nuclear radiation and liquid hydrogen

on mechanical properties of three phenolic materials', Advances in Cryogenic Engineering 11 (1 966) 486

113 Weleff, W. 'Final Report, GTR-1 6 radiation effects test on

structural materials at -423°F', Aerojet-General Corporation Report No RN-S-0290 (1966)

4. Handbooks and reviews

114 Coston, R. M. 'Handbook of thermal design data for multi- layer insulation systems', LMSC-A847882, Vol II (Final), George C. Marshall Space Flight Center, Huntsville, Alabama

(1967) (N67-34910)

1 1 5 Hertz, J. 'The effect of cryogenic temperatures on the mechanical properties pof reinforced plastic laminates' Society of Plastics Engineers Journal 21 (1965) 181

1 16 Hertz, J., Knowles, D. 'Survey of thermal properties of

selected materials'. General Dynamics/Convair Report AAL65-008 (AR-504-1-553) (1965) (N65-31775)

1 1 7 Jurevic, W. G., Rittenhouse, J. B. 'Structural applications handbook', AFML TR-67-332, Air Force Materials Laboratory,

Wright-Patterson Air Force Base, Ohio (1968) (AD 804 585)

1 1 8 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)

1 1 9 Landrock, A. H. 'Properties of plastics and related materials

at cryogenic temperatures', Plastic Report No 20, Plastics Technical Evaluation Center, Picatinny Arsenal, Dover, New

Jersey (1965) (AD 469 126)

1 20 Maximovich, M., Scheck, W. G. 'Data summary and reference

file for graphite and boron reinforced composite materials', General Dynamics Convair Report No GDCA-DBG7 1-006 (1971) (Contract NAS 8-26198, George C. Marshall Space

Flight Center, Huntsville, Alabama)

121 Nored, D. L., Hennings, G., Sinclair, D. H., Smith, G. T., Smolak, G. R., Stofan, A. J. 'Storage and handling of cryogenic fluids', NASA Special Publication SP-5053,

Proceedings of Conference on Selected Technology Cor the Petroleum Industry, Lewis Research Center, Cleveland Ohio

(1965) (N66-33674)

122 'Plastics for aerospace vehicles, Part 1 : Reinforced plastics', MIL-HDBK-17A, Department of Defense, Washington DC*

(1971)

123 Rittenhouse, J. B., Singletary, J. B. 'Space materials handbook',

3rd edn, NASA Special Publication SP-3051, National Aeronautics and Space Administration, Washington, DC(1969i

(limited publication as AFML-TR-68-205)

1 24 Schwartzberg, F. R., Herrzog, R. G., Osgood, S. H. et al 'Cryogenic materials data handbook (Revised), Volume II, AFML-TDR-64-280-Vol II (Revised), Air Force Materials

'

Laboratory, Wright-Patterson Air Force Base, Ohio (1970) (AD 713 620)

Miscellaneous references

125 Kastelic, J. R., Hiltner, A., Baer, E. 'Crazing, yielding and

fracture in polycarbonate and polyethelene terephthalate at low temperatures', Journal ofMacromolecular Science-Physics B7(4) (1973) 679

1 26 Relationships Between Structure and Mechanical Behavior in Polymeric Solids, ASM Materials Science Seminar, Chicago, 1 1 1, 1973 (in preparation)

CRYOGENICS. JUNE 1975 42

347
Bibliography — Property Cross Reference

Glass — vjiaSS — Glass — Glass — Glass — olass — silass —

_.

Property Epoxy ruiycoLcr Phenolic Teflon Silicone Pol yu re thane Phenyl Silane DDI** r d 1

1-1.3, 2-2.2, 1—1.3, 2—2.2, 1-1.3,2-2.2, 2-2.2,37,67, 1-1.3,2-2.2, 2—2.2 85 2—2.2, 57 67 2—2.2

3-3.5,4, 5, 3-3.2, 3.4, 40 54 55 57 68 80 94 24, 40, 57

5.1,22,24, 3.5, 34,40, 64—66 71 75 102 113 64 66 68 97

28, 33, 35 55, 57, 63, 83, 93, 97,

40, 43, 46, 64, 66, 75, 111, 112

47, 52, 54, 109

55, 57, 66,

71, 87

1-1.3,2-2.2, 1-1.3,2-2.2, 1-1.3,2-2.2, 2-2.2, 80 1-1.3,2-2.2, 2-2.2 2-2.2,57, 2-2.2, 24

3—3.5,4.5,5, 3,34,55,57 54,55,57, 24, 57 66

5.1,24,28, 97, 1 1

1

or rr cr oO, Ou, DO, 57,66,87 E\

1 — 1.3,2,3— 1-1.3, 2,55, 1.1.3,55,57 1-1.3,2,57 2 2, 57 2

O.O, ZO, OO, 57

56, 57

3-3.5,4,5, 3-3.5, 64 64,65,112 80,89,102 64 85

R 1 "}R fid O. 1 , OO, OH 85

a

fu 1-1.3,2-2.2, 1-1.3,2-2.2, 1—1.3,2—2.2, 2—2.2,80 1—1.3,2—2.2, 2-2.2 2.2.2 2-2.2

3-3.5,4,40, 3-3.3,3.5, 40,41,57,71, 40,57,97

57,66,71 df) R7 RR qo no 07

c 1 1—1.3,2, 1—1.3, 2, 1 — 1.3,2—2.2, 2—2.2,80 1 — 1.3,2—2.2, 2—2.2 2—2.2 Z—Z.z

3-3.2! 2.2, 57 R70 / Q7 , 00, so, y / R7

3.4, 3.5,57

2 1-1.3,2 1-1.3, 2 1 — 1.3, 2 1 — 1.3, 2 2 2

1 100 l— \.o,2., 1-1.3,2, 2.2, 1-1.3,2-2.2, 2-2.2,80 1-1.3,2-2.2, 2-2.2 2-2.2,57 2-2.2, 24

2..Z,6—3.L, 3-3.5,34,57 57,82,97 24,57

3.4 3.5

24, 57, 66 Ec 1,1.2, 1.3, 1-1.3,2,2.2 1-1.3, 2-2.2, 2-2.2, 80 1-1.3,2-2.2, 2-2.2 2-2.2 2-2.2

9 9 9 R7 57, 66 57,82,97 57

0

s1 3—3.2,3.4, 3-3.5

0. 0, H. 1 , 5,5.1,22,

47, 71

'i 1, 1.3,2,2.1, 1-1.3,2-2.2, 1-1.3, 2-2.2 1-1.3, 2 2 2 2

2.2,66,70, OH, /U 99

a

by 2-2.2,3,3.2, 2 2 2 2

3.4, 3.5 a

1 4,71, 103 62, 109 71,97 80 X 3,4.1,14,16- 3,14,16-16.3 14,16-16.3, 14,16-16.3 14,16-16.3, 14,16-16.3 16, 24

16.3,21,23- 90 49 24, 73, 90

25,38,4256

61,63,90,114

AL/L 3,5,14,16, 3, 14, 16-16.3 14, 16-16.3, 14, 16-16.3, 14,16-16.3, 5, 85, 91 14,16-16.3 16, 16.2 16.2,16.3,28, 26,27,50 67

69,85,91, 111 100,101,104, 114 CP 15,16,16.2 15,16-16.3 15,16-16.3, 15,16-16.3 15,16-16.3, 15,16-16.3, 16-16.3, 24

16.3,24,114 64 24

* includes fracture toughness

'* polybenzimidazole

348 CRYOGENICS . JUNE 1975 43
Biblopgraphy — Property Cross Reference (Cont'd)

Property

Graphite —

Epoxy

Boron —

Epoxy

Boron —

Aluminum

PRD-49

Epoxy

0

tu 8,9,10,10.2,

13,13.2, 32,

47, 51, 58

13, 20, 22, 47, 53, 58, 81

13, 13.3 29, 44, 59, 60

E\ 8,9,10,10.1,

13, 51

13, 81 13, 13.3 44, 59, 60

e« 13 20 13

8,10,10.1,

11, 12, 13, 13.2, 58

58

E

f 8, 10.1

,jCU 13 13, 13.3 £

c 13 13, 81 13, 13.3

8,9,10,10.1,

10.2, 12, 13, 13.2, 47, 58

13, 22, 47,

58, 81

13, 13.3 59, 60

45 20

13, 13.3 a

1 * 9, 10.2 81 X 13,17,18,

38

13,38,63,

81 AL/L 10,13,13.2,

17,18,36,51,

52,84,88

13 13, 13.3 59, 60

13 13 13, 13.3

* includes fracture toughness

Miscellaneous Composites

Glass-Polyimide o

tu (24, 67, 68), afu (89),. 0^(12.1, 89), AZ./M67)

Glass-Melamine (jtu (65), eM65), \(90)

Glass-Viton a

tu (68) Glass-Phenyl Formaldehyde 0*11 10), X (72), CL

(72) Si0 2 -Epoxy a

tu (31), AL/L (28)

Graphite-Aluminium o

tu (thermal cycling effects 6-6.2)

Graphite-Polyimide a

tu (13), a

fu (11, 12.1, 13), 0

s' (13), AL/L (36) Graphite-Phenolic X (14, 16-16.3), AL/L

(14, 16-16.3)

Steel-Aluminium a

tu

(7, 7.1), e

4 (7.1), a

1 (7, 7.1) Steel-Epoxy a

tu

(2, 30), £' (2, 30), a

fu (2), Ef (2), <f u (2),EC (2).n x (2) Boron/Steel-Aluminum c

,u

, E\ o

cu

, E

c

, Boron/Titanium-Aluminum a

by

, AL/L, Cp

, (13, 13.3), 0

s', e\ n* (13) Potassium Titanate-Epoxy A (16), Cp (16)

Miscellaneous Properties

Notch Tensile Strength Glass-Epoxy (3-3.2, 3.4, 3.5, 43, 46)

Vapor Permeability Glass-Epoxy (28) Modulus of Rigidity Glass-Epoxy (53), Glass-Teflon (80), Boron-Epoxy (53) Poissons Ratio Glass-Epoxy (53), Boron-Epoxy (53) Proportional Limit in Tension Glass-Epoxy (55, 56), Glass- Polyester (1-1.3, 55, 56)

Static Fatigue Glass-Epoxy (66, 99), PRD 49-Epoxy (44) Environmental Effects Glass-Epoxy (4-4.1), Graphite-Epoxy (13, 13.2, 17, 58) , Boron-Epoxy (13, 58), Boron-Aluminium (13, 36) PRD 49-Epoxy (44)

Electrical Resistivity Graphite-Epoxy (13), Boron-Epoxy (13) Thermo-Optical Effects Graphite-Epoxy (13, 84), Boron-Epoxy (13) Density Glass-Epoxy, Polyester, Phenolic, Silicone, Phenyl Silane

(14, 57), Glass-Teflon (14), Glass-Polybenzimidazole (16), Graphite-Epoxy (9, 13), Graphite-Phenolic (14), GraphitePolyimide (12.1) Radiation Effects (13.2, 33, 37, 40, 51, 52, 64, 65, 75, 94, 112, 113) Cryogen Compatability (33)

Pressure Vessel Applications

Glass-Filament

(19, 20, 39, 42, 46-48, 53, 67, 69, 76, 77, 91, 92, 95, 96, 98-100, 104-108, 113) Graph ite— Filamen t

(8, 10, 10.1,45,47, 100) Boron-Filament

(20, 47, 53, 78) PRD 49- Filament

(59, 60)

CRYOGENICS . JUNE 1975 349

44
Appendix IH : M. B. Kasen, "Properties of Filamentary-Reinforced Composites

at Cryogenic Temperatures/' Composite Reliability. ASTM STP

580, American Society for Testing and Materials, 1975 (to be

published)

.

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57
Appendix IV: R. E. Schramm and M. B. Kasen, "Static Tensile Properties of BoronAluminum and Boron-Epoxy Composites at Cryogenic Temperatures,"

Paper T3, Cryogenic Engineering Conference, Kingston, Ontario,

July 22-25, 1975 (to be published).

STATIC TENSILE PROPERTIES OF BORON-ALUMINUM AND

BORON-EPOXY COMPOSITES AT CRYOGENIC TEMPERATURES

R. E. Schramm and M. B. Kasen

Cryogenics Division

Institute for Basic Standards

National Bureau of Standards

Boulder, Colorado

Paper T-3

Preprinted for the

Cryogenic Engineering Conference

Kingston, Ontario, Canada

July 22-25, 1975

58
STATIC TENSILE PROPERTIES OF BORON-ALUMINUM AND

BORON-EPOXY COMPOSITES AT CRYOGENIC TEMPERATURES*

R. E. Schramm and M. B. Kasen

Cryogenics Division

Institute for Basic Standards

National Bureau of Standards

Boulder, Colorado 80302

ABSTRACT

State-of-the-art boron-aluminum and boron-epoxy composites have been

mechanically characterized at 295, 76, and 4 K. Static tensile properties

include elastic and shear moduli, Poisson's ratio, yield and ultimate strengths,

and elongation. The data are in a form useful in strength or stiffness limiting

predictions of complex crossply components using macromechanical composite

theory. Both composite types performed well at cryogenic temperatures.

Key Words: Advanced fiber composites; boron-aluminum; boron-epoxy;

cryogenics; static mechanical properties.

* This research was supported by the Advanced Research Projects Agency

of the Department of Defense under ARPA Order No. 2569; it is a contribution of the National Bureau of Standards, not subject to

copyright.

59
STATIC TENSILE PROPERTIES OF BORON-ALUMINUM AND

BORON-EPOXY COMPOSITES AT CRYOGENIC TEMPERATURES

R. E. Schramm and M. B. Kasen

INTRODUCTION

The Advanced Research Projects Agency is currently supporting research to

characterize the properties of various structural materials at liquid helium

temperature. The objective is to obtain design data for superconducting motors

and generators. Fiber-reinforced polymer-matrix composites are included in this

program because their low A/E and A/a ratios suggest the possibility of substantial savings in refrigeration cost compared to metal construction. An

aluminum-matrix composite is also included because the very high a/p and E/p

ratios of this electrically-conductive material suggests potential applications

in components such as the eddy-current damper shield.

A comprehensive survey [1,2] suggested that the boron-reinforced composites

require the least optimization for cryogenic service, as their room temperature

mechanical properties are relatively unaffected by cooling to 76 K. We therefore

selected boron-epoxy and boron-aluminum composites as the initial materials to

be characterized. The static tensile mechanical properties at 295, 76 and 4 K

of these materials are reported in this paper.

Macromechanical composite theories use key mechanical properties obtained

from tensile, compressive, and shear tests on uniaxial laminates to predict

strength and stiffness-limiting properties of complex crossply composites [3,4]

.

The key properties are the elastic constants, E^, E

22 > V^j G-^2' and the

associated strength parameters, a^ and Some analyses require additional

parameters, £-q> e22' anc* T12* T

^s notation refers to the composite axes defined in Fig. 1. A complete characterization requires evaluation of

compressive as well as tensile properties.

MATERIALS AND TEST METHODS

Table 1 lists the materials studied in this program. All panels were

obtained from commercial sources. Composite void content has not yet been

determined; however, radiography failed to reveal any voids. The average ply

60
thickness was 0.017 cm (0.0067 in) for each material. The composite panels

contained nominally 48% by volume 5.6 mil boron fiber. The 6061 aluminum

alloy matrix was in the F temper (as diffusion bonded) . The epoxy-matrix

composite was fabricated from commercial prepreg tape. Uniaxial longitudinal

(0°) tensile properties were determined from the 6-ply material, while the

15-ply material was used for evaluation of the uniaxial transverse (90°)

tensile properties. The + 45° panels were used to evaluate in-plane shear

properties. Ideally, shear properties would also be determined directly from

uniaxial laminates; however, the required specimens are expensive and require

special test fixtures. The + 45° tensile coupon method of calculating shear

moduli is a simplified test procedure [5,6],

All testing was performed in air (295 K) , liquid nitrogen (76 K) , or liquid

helium (4 K) on a universal testing machine at a strain rate of 0.01 min ^ using

a tensile cryostat of conventional design [7]. Load cell calibration was confirmed by dead weights. Resistance strain gages, parallel and normal

to the stress axis, were mounted on the uniaxial longitudinal and + 45°

specimens. The uniaxial transverse specimens were gaged only along the

tensile axis. Strain values were calculated from manufacturer's gage

factors and comparison with signal levels from precision shunt resistors

in the commercial bridges. The strain gages failed at about 1% strain at

cryogenic temperatures; however, larger strains were measured from crosshead

displacement (corrected for load train strain) and from displacement of gage

marks lightly scribed on the boron-aluminum specimens. Load and strain were

recorded simultaneously with a two-pen recorder.

Test specimens were straight-sided coupons, 2.5 cm (1 in) wide for transverse

and +45° orientations, and 1.3 cm (0.5 in) wide for the longitudinal fiber

orientation. The boron-aluminum specimens were 15 cm (6 in) long with a 5 cm (2 in)

gage length between grips; the boron-epoxy specimens were 28 cm (11 in) long with

a 15 cm (6 in) gage length. Specimens were cut with a diamond saw. Before

testing, the edges were lightly buffed with emery paper.

Grips were attached to the specimens in an alignment fixture. Specimen

ends, wrapped with 100 mesh stainless steel screen, were sandwiched between

serrated titanium plates attached to the load train by pins and clevises. Heavy

61
stainless steel pressure plates, each containing six pairs of screws, backed

the titanium plates and provided compressive force on the grip. The torque on

the screws was systematically increased toward the outer end of the grip to

uniformly transfer the tensile load into the composite specimen. When finally

assembled in the cryostat, the specimen alignment was within the limits prescribed in ASTM specification D 3039-71T [8].

The boron-epoxy specimens were conditioned for at least two days in 45 - 55%

relative humidity at room temperature before testing. The boron-aluminum did

not require preconditioning. At least three specimens were tested for each

material, orientation, and temperature. The boron-epoxy was initially cycled

three times within the elastic region (to about 25% of the tensile strength)

to establish the repeatability of the modulus, before increasing the load to

initiate fracture. The boron-aluminum was not cycled because of the low yield

strength of the aluminum matrix.

Tensile fractures in composite materials commonly occur at the specimen

grips, where a stress concentration inevitably exists. The gripping system

described here reduced the frequency of grip failures to about 20%. Tensile

strength data from specimens which failed close to the grips (within one specimen

thickness) were rejected, even though there was no clear evidence that such

failures affected the strength.

In-plane shear properties were calculated from analyses of the tensile

stress-strain curves of the + 45° specimens [5,6]. Poisson's ratio for the

transverse specimens, was calculate<i from the reciprocity relation [5]:

E

22 v = v 21 12 E n

where "*" S ^ rom uniaxial transverse specimens and and is from

uniaxial longitudinal tests.

RESULTS AND DISCUSSION

The static tensile properties and elastic constants for the boron-aluminum

and boron-epoxy composites are presented in Table 2 and 3. In-plane shear

moduli appear in Table 4. In accord with ASTM D 3039-71T, the number of

specimens, standard deviation, and coefficient of variation (standard deviation

62
expressed as a percentage of the average value) are included. For only

three or four tests, the calculated standard deviation is nearly equal to

the data spread. Figures 2 and 3 illustrate the temperature dependence of

the elastic moduli and tensile strengths of both composites.

Stress-strain curves for the uniaxial longitudinal test mode were

essentially linear to fracture for both materials at all test temperatures.

The uniaxial transverse boron-epoxy also failed with negligible yielding

at 76 K and 4 K; however, some plastic deformation was observed at 295 K.

All + 45° and the boron-aluminum uniaxial transverse specimens yielded

significantly prior to fracture at all test temperatures. The boron-aluminum

uniaxial transverse and + 45° specimens displayed discontinuous yielding

at 295 and 4 K, but not at 76 K.

Cryogenic temperatures had relatively little effect on the measured

mechanical properties. At 76 K, the averaged data indicate a minimum in the

elastic modulus of boron-aluminum for the transverse and + 45° orientations,

as well as a decrease in the calculated in-plane shear modulus. However, the

data scatter was such that this could be a statistical effect, rather than a true temperature effect. Significant temperature effects were noted in the

orientations where the matrix contributes to the properties. Thus, the

proportional limit of the + 45° boron-aluminum slightly increased on cooling,

while that of the boron-epoxy declined. Also, the elongation of the + 45°

boron-aluminum decreased sharply on cooling, while there was a relatively

small change for boron-epoxy.

The aluminum-matrix composite exhibited less anisotropy than the epoxymatrix composite, reflecting the higher strength and greater ductility of the

aluminum. The boron-aluminum has a transverse modulus approximately 75% of

that in the longitudinal direction, and about twice as high as that of the

unreinforced 6061 aluminum alloy.

The in-plane shear modulus values calculated for the boron-epoxy appear to

be in reasonable agreement with published room-temperature values [9]. However,

the shear moduli calculated for the boron-aluminum are inordinately low. It

has been shown that shear moduli calculated from analyses of + 45° tensile

stress-strain curves agree with moduli determined by more rigorous methods in

63
the case of graphite-epoxy and glass-epoxy composites [6]. However, Ramsey,

Waszczak, and Klouman [10] report that the nonlinear stress-strain response of

5.6 mil boron-6061 aluminum in the F condition may invalidate the simplified

+45° test method used here and lead to low shear moduli calculations for this

material. The boron-aluminum values here serve only as lower boundaries and

display the limits of this method.

ACKNOWLEDGMENT

The authors wish to thank Dr. Ralph Tobler of this laboratory for his

careful reading of this manuscript.

NOTATION

E = Engineering Elastic Modulus

G = Engineering shear modulus

e = Engineering strain

A = Thermal conductivity

V = Poisson's ratio

p = Mass density

a = Engineering stress

T = In-plane shear stress

64
REFERENCES

1. M. B. Kasen, "Properties of Filamentary-Reinforced Composites at

Cryogenic Temperatures," to be published in ASTM STP 580, Composite

Reliability .

2. M. B. Kasen, Cryogenics , 15(6): 327 (1975). (Also "Part II: Advanced

Composites" to be published in Cryogenics )

.

3. Advanced Composites Design Guide, Third Edition, Volume II, Analysis , Air Force Materials Laboratory, Wright-Patterson Air Force Base, Ohio

(1973).

4. M. E. Waddoups, in: Composite Materials Workshop , S. W. Tsai, J. C. Halpin, and N. J. Pagano, Eds., Technomic Publishing Co., Inc.,

Stanford, Ct. (1968), p. 254.

5. P. H. Pettit, in: ASTM STP 460, Composite Materials, Testing and Design , American Society for Testing and Materials, Philadelphia (1969), p. 83.

6. D. F. Sims and J. C. Halpin, in: ASTM STP 546, Composite Materials , Testing and Design , American Society for Testing and Materials,

Philadelphia (1974), p. 46.

7. R. P. Reed and R. L. Durcholz, in: Advances in Cryogenic Engineering , Vol. 15, Plenum Press, New York (1970), p. 109.

8. 1973 Annual Book of ASTM Standards, Part 25 , American Society for

Testing and Materials, Philadelphia (1973), p. 608.

9. Advanced Composite Design Guide, Third Edition, Volume IV, Materials , Air Force Materials Laboratory, Wright-Patterson Air Force Base, Ohio

(1973)

.

10. J. E. Ramsey, J. P. Waszczak, and F. L. Klouman, "An Investigation of

the Non-Linear Response of Metal-Matrix Composite Laminates," AIAA/ASME/

SAE 16th Structures, Structural Dynamics and Materials Conference,

Denver, Colorado (May 27, 1974).

65
LIST OF TABLES

Table 1. 5.6-mil Boron-Epoxy and Boron-Aluminum Composites.

Table 2. Tensile Properties of Boron-Aluminum Composites (Averages

of Specimens Tested)

.

Table 3. Tensile Properties of Boron-Epoxy Composites (Averages of

Specimens Tested)

.

Table 4. In-Plane Shear Modulus (Averages of Specimens Tested).

LIST OF FIGURES

Figure 1. Identification of composite Axes.

Figure 2. Temperature dependence of the elastic modulus of boron-epoxy

and boron-aluminum of various orientations. Error bars reflect

+ one standard deviation; in several instances, the error bars

are smaller than the datum symbol.

Figure 3. Temperature dependence of the ultimate tensile strength of

boron-epoxy and boron-aluminum of various orientations.

Error bars reflect + one standard deviation; in several

instances, the error bars are smaller than the datum symbol.

Table 1. 5.6-mil Boron-Epoxy and Boron-Aluminum Composites

Fiber Number of

Orientation Plies

Uniaxial 6

Uniaxial 15

+ 45° 10

66
S £

x> —

E 2

3 a

1

3

Si

5 =

.2 E w s

~ O ~H

(M (M f*l

^- CO CMN

00 O to

O O \Tl

-« 00 ^

2 to

5 3

« m (\j

r- rg

N o — o

h N N

o oc

>

s 1

,5

67
T3Vo

si si sL

68
Table 4 . In-Plane Shear Modulus3

(Averages of Specimens Tested)

Temperature

9 2 1(T N/m

Shear Modulus, G

10

6 psi

12

CV(%)

Boron-Aluminum

295

76

4

295

76

4

39.7(4.0)

27.0(5.5)

32.6(5.5)

Boron-Epoxy

4.72(0.44)

9.19(0.38)

9.28(0.19)

5.75(0.58)

3.92(0.80)

4.73(0.80)

0.684(0.064)

1.33(0.06)

1.35(0.03)

10.0

20.4

16.9

9.3

4.1

2.0

These values are calculated from cross-ply (+ 45°) data using the method

of Sims and Halpin [6], Standard deviations are in parenthesis; CV is

the coefficient of variation.

3The boron-aluminum moduli are only lower limits to the true values.

Fig. 1. Identification of composite axes,

59
(!sd s

oi] ^'Hi9Nms nvNinn

U fi 01 JJ w in 10 id

CO tH T3

01 O 0) H CO •- 4J

co 3 a

a o o a

CO (13 w

01 > -H

0 Q 0) t-t O C «

01 M O .£)

•O -O 4J ^

c c o u

0» M 01 01 a »h

01 >»«w 01 •a k 0) j=

0) a. M Ot CO » 3 I H CO u a a cu « o £> u

i-i u a

ot o u m

a- xi o u

; ui/N

8 0l] B^'m9N]HlS 31¥Wllin

(!sd

g oi)] sninaow ousvii

1^/HgOili sninooN susvu
NBSIR

SEMI-ANNUAL REPORT ON MATERIALS RESEARCH

IN SUPPORT OF SUPERCONDUCTING MACHINERY

ELASTIC PROPERTIES OF ENGINEERING MATERIALS

AT CRYOGENIC TEMPERATURES

H. M. Ledbetter, E. R. Naimon, and W. F. Weston

Cryogenics Division

Institute for Basic Standards

National Bureau of Standards

Boulder, Colorado 80302

October 1975

71
Summary: Elastic Properties

During the six months preceding October, 1975, the following studies were completed:

(1) Oxygen-free copper . Room-temperature elastic constants of this material were

determined by a pulse-echo method. Low-temperature Young's and shear moduli were

determined by a resonance method. Results of this study are included in the

accompanying manuscript describing copper-nickel alloys (See p. 73).

(2) Copper-nickel alloys . Two alloys, Cu-10 Ni and Cu-30 Ni, were studied as described

above for copper. Also, the longitudinal modulus of Cu-30 Ni was determined by a pulse-echo method at low temperatures. Advantages of combining resonance and pulseecho data for these materials are described along with results in the accompanying

manuscript "Low-temperature elastic properties of some copper-nickel alloys" by

H. M. Ledbetter and W. F. Weston. This study was presented at the 1975 IEEE Ultrasonics Symposium in Los Angeles, and it will be published in the proceedings of that

symposium (see p. 73).

'3) Copper alloy PD135 . This material, supplied and characterized by J. M. Wells of

Westinghouse, was studied at room temperature by a pulse-echo method to determine its

elastic constants. They are surprisingly different from those of unalloyed copper,

and they are described in an accompanying manuscript "Elastic properties of a coppercadmium-chromium precipitation-hardened alloy" by H. M. Ledbetter (See p. 77).

^4) Invar . Low-temperature elastic properties of an iron-36 nickel invar-type alloy were determined by both resonance and pulse-echo methods. This study is described in an accompanying manuscript "Low-temperature elastic properties of invar" by H. M. Ledbett

E. R. Naimon, and W. F. Weston. This study was presented at the 1975 ICMC-CEC in

Kingston, Ontario, and it will be published in the proceedings of that conference

(See p. 81).

5) Nickel-chromium-iron-alloys (Inconels) . There is an accompanying reprint "Dynamic

low-temperature elastic properties of two austenitic nickel-chromium-iron alloys,"

Mater. Sci. Engg. 20 (1975) 185-94 (See p. 95).

6) Austenitic stainless steels . There is an accompanying reprint "Low-temperature

elastic properties of four asutenitic stainless steels," J. Appl. Phys. 46_ (1975)

3855-60 (See p. 105).
LOW-TEMPERATURE ELASTIC PROPERTIES OF SOME COPPER-NICKEL ALLOYS*

H. M. Ledbetter and W. F. Westont

Cryogenics Division, National Bureau of Standards

Boulder, Colorado 80302

ABSTRACT. The polycrystalline elastic properties of Cu, Cu-10 Ni, and Cu-30 Ni were determined

between room temperature and liquid-helium temperature using both pulse-echo (10 MHz) and resonance

(60 kHz) methods. The temperature dependence of all three materials is regular. The composition

dependence of the elastic constants is reviewed and new composition dependences are proposed.

INTRODUCTION .. The study reported here was undertaken to elucidate the composition dependence and, especially, the

temperature depend-ence of the elastic properties of

copper-nickel alloys. Despite many previous studies,

these properties are not known satisfactorily.

Copper-nickel alloys are well known for their re- sistances to corrosion and to stress-corrosion cracking.

They are used in condensers and in heat exchangers.

'Recently, these alloys have become candidates for applications in superconducting machinery.1 The elastic properties (Young's modulus, shear modulus, etc.) of these

alloys have been studied extensively at .room temperature,. and one high-temperature elastic-property study

jhas been reported. At cryogenic temperatures only the

Young's modulus has been studied, and only at liquidnitrogen and liquid-helium temperatures. The complete

I set of polycrystalline elastic properties is reported

here for Cu, Cu-10 Ni, and Cu-30 Ni semi-continuously

between room temperature and liquid-helium temperature.

The elastic constants are essential engineering design

parameters, and they are related simply to interatomic

forces. Copper-nickel alloys are especially interesting elastically because nickel increases copper's

elastic stiffness, contrary to usual alloying effects.

Materials

Materials were obtained from commercial sources in the form of 1.9-cm (3/4-in) rods. Their chemical analyses by weight are: 9.98 Ni, 1.16 Fe, 0.07 Zn, 0.02' Pb

for the Cu-10 alloy; 30.05 Ni, 0.72 Mn, 0.59 Fe, 0.04

Zn, <0.01 Pb, <0.01 Sn, <0. 01 Al, <0.01 As. The copper

-was an oxygen-free high-conductivity grade. Their physical and metallurgical characterizations are given in Table I.

Table I. Physical and metallurgical

characterization of the materials studied

Mass density (g/cm 3

) Hardness (DPH No.

1 kg load)

Average grain size (mm)

by linear intercept

method

Cu 8.952

45

1.0

Cu-10 Ni

8.903

87

0.05

C u-_3 0_ Ni

8.914

124

0.03

Condition Cold drawn

60%,

annealed

Annealed

677 °C, 40 min

Commercially

annealed

Experimenta

Two methods, resonance (60 kHz) and pulse-echo

(10 MHz) , were used to determine sound velocities in the

circular cylinders 0.48 cm (3/16 in) in diameter and

about 3 cm and 2 cm long for the longitudinal and transverse cases, respectively. The pulse-echo speci- mens were 1 cm cubes. Temperatures were monitored with

chromel-constantan thermocouple's

The problem of the effect of texture on the elastic constants was studied by measuring the sound

velocities both parallel and perpendicular to the

rolling axis and also measuring the shear-wave birefringence for the latter case. The shear and longitudinal wave velocities were found to vary with direction by an average of 0.6 and 0.3 percent, respectively. The -birefringence in the shear velocity was

found to have a maximum value of 1.5 percent. Thus,

the specimens had negligible texture, and the velocities were simply averaged over direction.

Results

Some of the low-temperature elastic constants are shown in Figs. 1-3, and room-temperature values are given in Table II. Young's modulus E was determined by

a resonance method, the shear modulus G by resonance

(shown in Fig. 2) and by pulse method, and C ( by pulse

method

Table II. Elastic constants of Cu-Ni alloys at room temperature, units of 10 1

' N/r

( dimension less)

except \) h

B E G V Cu 2. 013 1 376 1 286 0 478 0 345

Cu-lONi 2. 092 1 424 1 345 0 501 0 343

Cu-30Ni 2. 250 1 517 1 472 0 550 0 339

While there are only two independent elastic con- stants for an isotropic material, several are useful

for various applications. Besides E, G, and C^ already mentioned, the bulk modulus B' (reciprocal com- pressibility) and Poisson's ratio v are often required.

Relationships among these elastic constants and their

relationship to the sound velocities are well known:

and

G = pv , t

c

i = pv r

B = C

£ "

3

•F = 9BG

3B + G

E

V = 2G -]

(1)

(2)

(3)

(4)

(5)

Since three elastic constants were measured — E, alloys. Both methods were described previously.' 3 Low- and C 0 — the problem is, in principle, overdetermined

temperature pulse-echo data were determined by a super- position method. The resonance specimens were right

* Contribution of NBS, not subject to copyright. Present address: Rocky Flats Plant, Rockwell Inter- national, Golden, Colorado, 80401.

However, as can be inferred from Table III, this is effectively not the case. It is to be recalled that in

a pulse-echo experiment V£ and v

t are measured with

typical errors of + 1/2 percent. In a resonance experiment, E and G are measured with tvpical errors of + 1 percent
Table III. Maximum percent propagation errors

involved in elastic-constant determinations,

assuming v = 1/3. Boxed parameters are

. those measured. j2_Lb.exs__.ar_e-- V

i V

t C

l

E G V B

Pulse-echo 1/2 1/2 1 1 1 2/3 1 Resonance 5-1/2. 1/2 11 1 1 8 17

Thus, when E and G are determined (as in a resonance

experiment), derived values of C^, V, and B are quite

inaccurate compared to pulse-echo results. Usual re- sonance methods are useful for determining E and G, but much less useful for determining the other elastic

constants. The superiority of the pulse-echo method

is apparent. However, in many cases, including copper- nickel alloys longitudinal ultrasonic velocities can

be determined accurately while transverse ultrasonic

velocities cannot, because of higher attenuation.

Thus, it becomes advantageous to combine pulse-echo

data (C^,) with resonance data (E or G) . In the present work, it was found that the resonance value of E was slightly more accurate than the resonance value of

G. Thus, the C^, (pulse-echo) - E (resonance) pair is the best choice. Then it is necessary to solve a quadratic equation for G:

2 .1/2

(E + 3C 0 ) + [ (E + 3C 0 ) 16EC,

G = (6)

discarding the non-physical root. The data in Figs 1-3 were fitted to a theoretical

relationship suggested by Varshni'1 :

C

s

t?~T

1 .(7)

where C , s, and t are adjustable parameters, and T

is the temperature.

The composition dependence of the elastic con- stants is shown in Figs 4-9, where the present results

are compared with results of sixteen previous stu- dies? 20 Two of these previous studies were made on single crystals rather than on polycrystalline aggregates. These single-crystal results were averaged by a Voigt-Reuss-Hill method.21 The lines in Figs 4-8 are linear least-squares fits to the averaged singlecrystal results and to the present results.

Temperature coefficients of the elastic constants

are given in Table IV.

2. The composition dependence of the elastic con- stants is approximately linear over the composition range studied. The Young ' s-modulus/

composition curve goes through the value for

copper (1.27 x 10 1 N/m 2

) and not through the

fictitious value of E for Cu (1.30 x 10 11 N/m2 ),

2 2 as suggested by Koster. The parameters pro- posed here for the change of E and G with com- position (AC/Ax = 0.690 and 0.259, respectively,

where C denotes elastic constant and x denotes

atomic fraction) differ from those proposed by

Subrahmanyan 2

0

(0.798 and 0.319, respectively).

3. In some cases, there is considerable advantage

in studying elastic properties of solids by

both pulse and resonance methods. In the presenl case, the specimens transmitted longitudinal

waves thus permitting the determination of C^. But they damped shear waves, precluding determination of G by a pulse method. G, or alternatively E, could be determined by a resonance

method. The C„-E data gave better values of B and vj than could be obtained from resonance experiments alone. t

Acknowledgment •

This work was supported by the Advanced Research

Projects Agency of the U.S. Department of Defense.

References

1. .Handbook on Materials for Siiper conducting

Machinery, Battelle Metals and Ceramics Information Center, Columbus, Ohio, 1974.

2. E. R. Naimon, W. F. Weston and H. M. Ledbetter,

Cryogenics 14 (1974) 246-9.

3. W. F. Weston, J. Appl. Phys., forthcoming.

4: Y . P. Varshni, Phys. Rev. B 2 (1970) 3952-8.

5. E. Gruneisen, Ann. Phys. Leipz. [IV] 22 (1907)

, 801.

6. Z. Nishiyama; Sci. Rep. Tohoku Univ. Sendai

18 (1929) 359-400.

7. S. Aoyama and T. Fukuroi, Bull. Sci. Res. Inst. Tokyo 20 (1941) 809-11. In Japanese.

Table IV. Temperature derivatives of elastic constants at room temperature (10 ""K

1 )

Alloy 1 dB

B dT

1 dE

E dT

1 dG • G dT

1 dv

V dT 10

Cu -1.64 -3. 64 -4. 01 4.65

Cu-lONi -3. 78 -3.61 11

Cu-30Ni -1.29 -2. 30 -3.20 8.22

12 Discussion and Conclusions

1. The low-temperature elastic properties of Cu-Ni

alloys resemble those of copper, except for scaling factors. Like Cu, their behavior is quite regular—monotonically decreasing, linear

at high temperature, lower slope at lower tem- peratures, and zero slope at zero temperature.

13.

14.

H. Masumoto and H. Saito, Nippon Kingoku GakkaiShi 8 (1944) 49-55. In Japanese.

T. Fukuroi and Y. Shibuya, Sci. Rep. Tohoku

Univ. A 2 (1950) 748.

W. F. Simmons, B. J. Sirois, D. N. Williams and

R. J. Jaffee, ASTM Proc. 59 (1959) 1035-45.

R.' E. Schmunk and C. S. Smith, Acta Metall. 8 (I960) 396-401.

B. Subrahmanyam and Bh. Kr ishramur ty , Indian J. Pure Appl. Phys. 1 (1963) 185-7.

J. B. Greer and E. H. .Bucknall, Trans. ASM 57 (1964) 555-60.

A. F. Orlov and S. G. Fedotov, Fiz. Metal.

Metalloved. 22 (1966) 137-8.
R. P. Reed and R. P. Mikesell, J. Mater. ASTM

2 (1967) 370-92.

G. Faninger, Z. Metallk. 60 (1969) 601-5.

K. Salama and G. A. Alers, IEEE Trans. Sonics

Ultrasonics (Jan. 1969) 28. Abstract.

L. M. T. Hopkin, H. Pursey and M. F. Markham,

Z. Metallk. 61 (1970) 535-40.

H. Masumoto, H. Saito and S. Sawaya, Trans. Japan Inst. Met. 11 (1970) 88-90.

B. Subrahmanyam, Mater. Sci. Eng. 15 (1974)

177-9.

R. Hill, Proc. Phys. Soc. Lond. A65 (1952) 349-54.

W. Koster, Mater. Sci. Eng. 18 (1975) 289.

0.620 0600 cz 0.580 o

to 0.560

r>

_j3 0.540 a

o

0.520

or< 0.500

X

CO 0.480

0460

"i r

Cu- i-.ONi

-a>raXD:*£[f~€K> Cu-10 Ni

Fig.

0 50 100 150 200 250 300

TEMPERATURE, K

2, Shear ifiodulus of tonic Cu-Ni alloys vcr^air; temperature. Resonance results.

CM

CO3

_l

=1Q

O

CO "o3

o>-

.700

1.600

.500

1.400

.300

.200

1 1 1 1 1

Cu-30 Ni

Lev i i III"' cir=5T-~£H~'0Ni

1

1

1 1

1 0 50 100 150 200 250 300

TEMPERATURE, K

Fig. 1. Young's modulus of sorr.e Cu-Ni alloys versus

temperature. Error bars are + 2a where a is

the standard deviation of nine measurements

and they reflect both experimental errors

and material variability; 2a represents, approximately, the 957. confidence interval.

P.esonance results. .

Fig. 4. Bulk modulus and Poisson ratio of Cu-30 Ni,

G (resonance)

Fig.

50 100 150 200 250 300

TEMPERATURE, K

3. Longitudinal modulus versus temperature

for a Cu-30 Ni alloy. Pulse-echo results.

PJ C 1.570 -

CO

ZD_l

ZDO

o

00

I 550 -

530

1.510 50 100 150 2 00 250 3C0

TEMPERATURE, K
LEGEND

• Grunoison ( 1907} e Nishiyomo (1929)

a Aoyamo 8 Fufcurr* (1941) « Mcsurrioto a Saito (1944)

<D Fukuroi S Shibuya (1950) C Simmons et ol ( 1959)

0) Schmun* a Smith (I960)

t» Greer a Buchnoll (1964)

^ Fantnqsr (1969) > Solama 5 Alers (1969 , unpublished

)

O Hophln et ol, (1970) < Moaumoto et ol (1970)

to Subrahmanyam a Krishnamurty (1963) o Subrohmonyom (1974) • Present Study

NICKEL (at. pet.)

Fig. 5. Compositional variation of for Cu-Mi

alloys. Half-filled diamonds and circles

represent single-crystal results averaged

by the Voigt-Reuss-Hill arithmetic method.

Line is a least-squares "it of present

results and averaged single-crystal results.

See first figure caption for meaning of error bars.

10 20 30

NICKEL (at. pet.)

Fig. 7. Compositional variation -of G for Cu-Ni alloys.

10 20 30 40

NICKEL (at. pet.)

Fig. 8. Compositional variation of B for Cu-Ni alloys.

i I i.l

Fig. 6.

10 20 30 40

NICKEL (at. pet.)

Compositional variation of E for Cu-Ni alloys.

<

oro

CO

CO

.38

.36

-1 1 r-

.34

2.32

.30 i l i

t 1 r r

j L 10 20 30 40

NICKEL (at. pet.)

Fig. 9. Compositional variation of V for Cu-Ni alloys.

76
ELASTIC PROPERTIES OF A COPPER-CADMIUM-CHROMIUM PRECIPITATION-HARDENED ALLOY*

H. M. Ledbetter

Cryogenics Division

National Bureau of Standards

Boulder, Colorado 80302

Abstract

The room-temperature elastic constants of a polycrystalline

precipitation-hardened copper-cadmium-chromium alloy were determined

by a 10 MHz pulse-echo method. With respect to copper the alloy has

a fifteen-percent-lower bulk modulus, a fifteen-percent-higher shear

modulus, and a four teen-percent-lower Poisson ratio. These changes,

especially in the Poisson ratio, are much larger than those observed

in the more familiar copper-beryllium precipitation-hardened alloy

and compare to those usually obtained only by mechanical deformation

or by phase transformation.

Key words: Bulk modulus; compressibility; copper alloy; elastic

constant; Poisson ratio; precipitation hardening; shear modulus; sound

velocity; Young's modulus.

The elastic constants of a material are essential in engineering design for predicting

load-deflection behavior. They are also essential for understanding material behavior related to interatomic bonding forces. Effects of many metallurgical variables on elastic

constants have been studied extensively* 3 these variables include alloying, annealing,

magnetic field, mechanical deformation, phase transformation, pressure, stress, and temperature. This subject was reviewed recently for copper by Ledbetter and Naimon^

Some experimental results on the elastic constants of a precipitation-hardened copper

alloy are described here. Compared to previous studies, the observed

elastic-constant changes are large.

The studied materialt is characterized as follows: Chemical composition by weight is

0.31 Cd, 0.33 Cr, < 0.01 Pb, < 0.01 Fe, < 0.01 Ni, < 0.01 Mn, balance Cu. Mass density

determined by hydrostatic weighing is 8.94 g/cm 3. Hardness is 57.5 on a Rockwell 30-T

scale. ASTM grain size number is five. The material was produced by extruding a 20.3-cm

diameter, 40.6-cm long billet at 1227 K and following with an 866-K, 1-hour precipitation

heat treatment.

The elastic constants were determined by measuring the ultrasonic longitudinal-wave

and shear-wave velocities, v^ and v , using a McSkimin 5 pulse-echo method near 10 MHz.

The usual engineering elastic constants are related to these velocities by: 6

2 Shear modulus = G = pv , (1)

2 Longitudinal modulus = C^ = Pv^, (2)

Bulk modulus = B = - j G, (3)

Young's modulus = E = 3GB/ (C^ - G) , (4)

and Poisson ratio = v = (E/2G) - 1 = \

(C

£ - 2G)/(C

£ - G) , (5)

where p is the mass density.

* Contribution of NBS, not subject to copyright.

This material was obtained from the Phelps-Dodge Corporation, denoted as their alloy

PD135. This trade name is used here to characterize the material; it is not an NBS

endorsement of this material.

77
The specimen consisted of a 0.95 x 1.35 x 1.35 cm "cube" whose opposite sides were

flat and parallel within 2.5 ym. For each of the three cube directions, longitudinal-wave

and shear-wave sound velocities were measured. Shear-wave birefringence was found to be

negligible. Both longitudinal and shear velocities varied less than one percent with

direction. Thus, velocities were simply averaged over the three directions.

Results of the present study are given in Table I. For comparison, similar data are shown for unalloyed copper and for a copper -1.85 beryllium precipitation-hardening alloy.

Copper data are from the single-crystal results of Overton and Gaffney 7 averaged by a Voigt-Reuss-Hill-arithmetic method? Such data agree with data obtained by measuring

polycrystalline specimens, and they are believed to be more reliable^ Cu-Be data are due

to Frederick and were reported by Richards? They were obtained from sound-velocity mea- surements on a 1.4-cm diameter specimen obtained from material that was cold drawn and

then precipitation hardened.

Table I. Room- temperature elastic constants of copper, precipitation-hardened

Cu-1.85 Be, and precipitation-hardened Cu-0.3 Cd-0.3 Cr. Velocities

have units o f 10 6 cm/ sec; elastic constants have units o f 10 11 N/m2

; v is dimensionless . Quantities in parentheses were calculated from

reported results. Error limits represent one standard deviation; they

reflect both experimental errors and elastic anisotropies

.

Cu Cu-Be Cu-Cd-Cr

p 8.93 8.332 8.936

v

i 0.4733 (0.4848) 0.4582 + 0.0053

v

t 0.2295 (0.2414) 0.2459 + 0.0021

2.000 (1.958) 1.876 + 0.043

G 0.473 0.485 0.540 + 0.009

B 1.371 (1.311) 1.156 + 0.016

E 1.271 1.296 1.402 + 0.012

V 0. 346 0.335 0.298 + 0.003

Results in Table I show for the copper alloys considered that the effect of precipitation hardening is to decrease C^, B, and v and to increase G and E. Since there are only

two independent elastic constants for a polycrystalline (isotropic) aggregate, these changes

in the elastic constants are not independent; and it is useful to consider their interrelationships. The most fundamental choice of an elastic-constant pair is the pair B and G

because these constants characterize the two extreme types of mechanical deformation —

volume change without shape change and shape change without volume change, respectively.

From Eqs. (l)-(4) it follows that:

dC

£

=

B dB 4G dG

C

£

" B +f G

B 3(B+| G) G

dE

E

(6)

(7)

dv 9GB , dB dG . , a

, and ~ =

(G + 3B)(3B - 2G)'

( eT " G~ }

' (8)

For a typical value of v, near 1/3, these relationships become:

dC_! 2 dB 1 dG (6a)

C

£ ~ 3 B 3 G

78
Thus, it is clear that E and G will change similarly with respect to most metallurgical

variables, that C„ depends more on B than on G, and that v is expected to be relativelyinvariant.

There are three principal unexpected results of the present study. First, the elastic

constants are changed significantly — G is increased fourteen percent and B is decreased

fifteen percent with respect to unalloyed copper. Second, Poisson's ratio is changed significantly. Koster and Franz showed that V is relatively insensitive to most metallurgical

variables except mechanical deformation and phase transformation. Both these variables

introduce strong elastic anisotropies , and it is not clear that these anisotropies have

been carefully considered with respect to changes in v. In the present case, however, no elastic anisotropy due to precipitation is expected and none was detected experimentally.

Thus, present results demonstrate a large change in the elastic constants of copper due to precipitation. As shown by Eq. (8a) , the unusually large change in V is due to B and G

being oppositely affected by precipitation. Third, the opposite behavior of B and G is

unexpected; usually G and B are roughly proportional, typically G ~ 3/8 Bi° As shown by

the data in table I, G/B is 0.35, 0.37, and 0.47 for copper, copper-beryllium, and coppercadmium-chromium, respectively.

Understanding the non-parallel behavior of B and G due to precipitation hardening

emerges as the principal problem of the present study. Very few experimental studies on

the elastic properties of precipitation-hardening alloys have been reported. The usual

elastic theories of precipitation 11 cannot explain the anomaly, either by precipitatematrix, precipitate-precipitate, or combined elastic-interaction energies. Also, the

problem remains to assure that this anomaly is not an artifact, an effective slowing of the

longitudinal ultrasonic wave due to its interactions with the precipitates. If the apparent

longitudinal velocity is real, then it may be necessary to invoke two mechanisms to explain

the elastic-constant changes due to precipitation hardening. One mechanism would increase

(decrease) both B and G while the other would decrease B (increase G) , leaving G(B) unaffected .

Acknowledgment

This work was supported by the Advanced Research Projects Agency of the U.S. Department

of Defense. Dr. J. M. Wells of Westinghouse supplied the material and its characterization.

References

1. R. F. S. Hearmon, Rev. Mod. Phys. 18 (1946) 409, Adv. Phys. 5 (1956) 323.

2. H. B. Huntington, in Solid State Physics, vol. 7 (F. Seitz and D. Turnbull, Eds.),

Academic, New York, 1958.

3. W. Koster and H. Franz, Metall. Rev. 6 (1961) 1.

4. H. M. Ledbetter and E. R. Naimon, J. Phys. Chem. Ref. Data 3 (1974) 897.

5. H. J. McSkimin, J. Acoust. Soc. Amer. 33 (1961) 12.

6. G. Bradfield, in Physical Examination of Metals (B. Chalmers and A. G. Quarrell, Eds.),

Arnold, London, 1960, p. 565.

7. W. C. Overton, Jr. and J. Gaffney, Phys. Rev. 98 (1955) 969.

8. R. Hill, Proc. Phys. Soc. Lond. A65 (1952) 349.

79
9. J. T. Richards, in Symposium on Determination of Elastic Constants, Spec. Tech.

Publ. No. 129, ASTM, Philadelphia, 1952, p. 71.

10. C. Zwikker, Physical Properties of Solid Materials, Pergamon, London, 1954, p. 90.

11. G. Sauthoff, Z. Metallkde. 66 (1975) 106.

80
LOW-TEMPERATURE ELASTIC PROPERTIES OF INVARtf

H. M. Ledbetter, E. R. Naimon,*t and W. F. Weston*

Cryogenics Division

Institute for Basic Standards

National Bureau of Standards

Boulder, Colorado 80302

ABSTRACT

The elastic properties of a polycrystal 1 i ne invar alloy were determined

between room temperature and liquid-helium temperature by two methods: mea- surement of ultrasonic (10 MHz) wave velocities with a pulse-echo technique,

and measurement of resonance frequencies (60 kHz) of cylindrical specimens

with a composite piezoelectric-oscillator technique. The shear moduli obtained

by the two methods are essentially the same. However, the other elastic con- stants (all of which have a dilatational component) differ both in magnitude

and in temperature dependence. Present pulse-echo results agree closely with

previous ' results obtained for both polycrystals and single crystals in a saturating magnetic field. The following elastic constants are reported:

longitudinal modulus, Young's modulus, the shear modulus, the bulk modulus

(reciprocal compressibility), and Poisson's ratio. The role of magnetic

effects on invar's elastic properties is discussed briefly.

Key words: Bulk modulus; compressibility; Debye temperature; elastic con- stant; invar; iron alloy; nickel alloy; Poisson's ratio; shear

modulus; sound velocity; Young's modulus.

ft Contribution of NBS, not subject to copyright.

* NRC-NBS Postdoctoral Research Associate, 1973-4,

t Present address: Dow Chemical, Rocky Flats Division, Golden,

Colorado 80401

.

81
LOW-TEMPERATURE ELASTIC PROPERTIES OF INVARtt

H. M. Ledbetter, E. R. Naimon,*f and W. F. Weston*

Cryogenics Division

Institute for Basic Standards

National Bureau of Standards

Boulder, Colorado 80302

INTRODUCTION

Because invar has nearly temperature-invariant dimensions, it is studied

scientifically and used technologically [1-5]. It is a face-centered-cubic

iron-base alloy containing about thirty-six percent nickel. Invar's unusually

low thermal expansivity is related to its magnetic properties [3]. The usual

thermal expansion is canceled approximately by magnetostriction, the coupling

between magnetization and strain.

Other physical properties of invar are also anomalous. It has positive

thermoelastic coefficients over a large temperature range. This anomaly is

useful in developing alloys with nearly temperature-invariant elastic constants.

Knowledge of invar's elastic constants is important. First, elastic

constants are related simply to interatomic potentials, and they reflect the

nature of the interatomic bonding. Second, they are essential design parameters

for predicting deflections in stressed components; they become particularly

important in cases of large stresses, large deflections, or acute tolerances.

The elastic constants of invar are of interest at low temperatures because

invar is a candidate low-temperature material [6], and (as discussed below)

its elastic behavior is especially anomalous below about 50 K.

The elastic properties of invar between room temperature and liquidhelium temperature are reported in the present paper. Elastic constants

were determined by two dynamic methods. Longitudinal and transverse ultrasonic wave velocities were measured at 10 MHz using a pulse-echo technique.

And the resonance frequencies of both longitudinal and transverse 60 kHz

tt Contribution of NBS, not subject to copyright.

* NRC-NBS Postdoctoral Research Associate, 1973-4.

t Present address: Dow Chemical, Rocky Flats Division, Golden,

Colorado 80401.

82
standing waves were determined by a three-component piezoelectric compositeoscillator technique. The following elastic constants are reported: Young's

modulus E, the shear modulus 6, the bulk modulus B (which is the reciprocal

compressibility), Poisson's ratio v, and the longitudinal modulus C^.

PREVIOUS STUDIES

Five experimental studies on the low- temperature elastic properties of

invar have been reported [7-11], and these studies are summarized in Table I. Despite considerable experimental study on them, the low-temperature elastic

properties of invar are not well characterized. The curves of Durham et al . [7] are based on a few fixed-temperature measurements, the lowest of these being

at 20 K. The extrapolation of these data to lower temperatures is questionable.

The Meincke and Litva [8] data are limited to temperatures below 160 K. Fletcher [9] reported only relative values and limited his measurements to

temperatures below 77 K. Maeda's [10] values are about twenty percent lower

than other reported results for E and G. Hausch and Warlimont [11,12]

measured single-crystal elastic constants; the averaging process for obtaining

polycrystal 1 ine elastic constants from single-crystal elastic constants is

uncertain; in particular, averaging has not been discussed for magnetic

materials that may have a magnetostrictive strain component. Also, the

Hausch and Warlimont data do not extend below 77 K. Thus, further experimental

studies are required to better characterize the low- temperature elastic properties of invar.

High-temperature elastic-property studies were reported by: Chevenard [13];

Kbster [14]; Fine and Ellis [15]; Hill, Shimmin, and Wilcox [16]; and Tino and

Maeda [17].

Besides dependence on variables such as temperature, mechanical deformation,

and magnetic field, the elastic properties of invar are sensitive to nickel

content. This variable must be considered carefully in comparing results from

specimens of different materials, particularly near thirty-six percent nickel

where properties such as E and G have sharp minima [18]. .

83
EXPERIMENTAL

For brevity, most experimental details are omitted here. Both the

pulse-echo technique [19] and the resonance technique [20] are described

elsewhere.

Samples of invar were obtained from a commercial source in the form

of 3/4-in (1.9-cm) diameter rods. The room- temperature mass density of the

3 sample was found to be 8.084 g/cm by an Archimedean method. The diamondpyramid hardness number was 226 for a 1 kg load. Grain size was found by

a linear intercept method to be approximately 0.02 mm. The composition of

the sample, obtained from a mill analysis, was: 0.08 C, 0.01 P, 0.01 S,

35.99 Ni, 0.81 Mn, 0.35 Si, 0.17 Se, and the balance Fe. Materials were

tested in the as-received condition, about fifteen-percent cold drawn.

Annealing the specimens increased their attenuations, reducing considerably

the quality of their pulse-echo patterns. (Magnetic domain walls attenuate

more in the annealed state than in the mechanically deformed state.)

The possible existence of preferred orientations (textures) in the

specimen was investigated by measuring the longitudinal sound-wave velocity

v^ in three orthogonal directions at room temperature. The longitudinal

modulus is given by:

where p is the mass density. Thus, both dilatational and shear deformation modes are sampled simultaneously. The results for the three directions

are: 0.4821, 0.4825, and 0.4929 x 10

6 cm/sec for the two transverse and for

the longitudinal axes, respectively (longitudinal axis = rolling direction).

These results were interpreted to indicate that the specimen had a slight

preferred orientation, but that practically it could be ignored. The rolling

direction of the bar was chosen as the wave-propagation direction for the

pulse-echo experiments, and it was the specimen axis for the resonance

experiments.

84
RESULTS

Results of the present study are shown in Figs. 1-5, together with the

low- temperature elastic constants of invar reported by others. Values of

the elastic constants at selected temperatures are given in Table II. Based

on previous measurements on standard materials, errors in E, G, B, and

determined by a pulse-echo method are estimated to be ± 1%. Errors in E and

G by a resonance method are also estimated to be ± VL Errors in the pulseecho value of v are larger. In particular, values of B, C^, and v obtained

from resonance data tend to be inaccurate because of the propagation of mea- surement errors; these derived resonance values are not reported here. For

the pulse-echo case, the elastic constants were obtained from the following

formulas:

G = pv^ , (2)

B = C

£

- | G , (3)

E = 9GB

(4)

G + 3B '

K

'

and v =

|g- - 1 , (5)

where v^ is the transverse wave velocity.

DISCUSSION

Pulse-echo Results

The present pulse-echo results agree within a few percent with results

reported by Meincke and Litva [8] for polycrystal 1 ine specimens, and also

with the results obtained by averaging the single-crystal results of Hausch

and Warlimont [11] by a Voigt-Reuss-Hi 1 1 method. Both these experiments

were done in a saturating magnetic field; the present experiments were done

in. zero applied magnetic field.

The agreement between the zero-field and the saturated-field results

can be understood as follows. At a high frequency (10 MHz in this case),

85
domain-wall motion is suppressed since the domain walls cannot follow the

high-frequency applied stress. In nickel, it was reported that the decrement

is a maximum at 150 kHz [21]; thus, domain-wall motion can contribute

significantly to the strain at this frequency. At higher frequencies the

domain-wall motion contributes a smaller strain. Therefore, the elastic

stiffnesses determined at high frequencies for unmagnetized materials are nearly as large as the elastic stiffnesses determined for magnetically saturated

materials. Thus, present pulse-echo results are consistent with the previous

three pulse-echo experiments [8,9,11], independent of magnetic field.

The present pulse-echo experiments were precluded below 40 K because

of the high attenuation of elastic waves in invar at these temperatures.

This attenuation phenomenon was detected first by Meincke and Litva [8], and

it was studied in detail by Fletcher [9] who described it by a general relaxationmechanism.

Resonance Results

The present resonance results disagree with the results reported by

Maeda [10], both in magnitude and in temperature dependence. Present results

are about twenty percent larger for both E and G, and the present temperature

coefficients are higher. It is difficult to understand this discrepancy.

It may be a frequency-dependent effect. If so, it is a large effect and

deserves further study.

Pulse-echo and Resonance Results Compared

For the shear modulus, present resonance and pulse-echo results are

identical within experimental error. The unusual behavior of the shear

modulus of invar is shown in Fig. 1. A regularly behaved material has a shear modulus that decreases monotonical ly with increasing temperature.

The two temperature coefficients are approximately the same, although the

slope in the resonance case is slightly higher than in the pulse-echo case. The resonance data for G also show a minimum near 50 K and a slight upturn

at lower temperatures, similar to the previously reported pulse-echo data [8-9].

86

i,'
For Young's modulus, the resonance and pulse-echo values differ by

a few percent in magnitude, but considerably in their temperature dependence.

These differences are shown in Fig. 2. The resonance value of E reaches

a minimum near 150 K and then increases steadily with lower temperatures. If

this is related to the minimum of E near 50 K shown by the pulse-echo data,

then a frequency-dependent attenuation mechanism is indicated. Fletcher [9]

found no appreciable frequency dependence of the damping in his experiments

between 10 and 35 MHz, but these frequencies may be so high that the attenuation

mechanism associated with the minimum is immeasurably small. It is interesting

that the resonance values of E are higher at all temperatures than the pulseecho values. From the usual AE^ effect, an elastic softening would be expected

in going from pulse-echo to resonance values [5]. Maeda [10] did find an elastic stiffening due to the AE^ effect, as evidenced by increased values

of E in a saturating magnetic field.

From the results shown for G and E, the disparity between resonance

values and pulse-echo values of the elastic constants would be expected for

any elastic constant that contained a dilatational component. Besides

Young's modulus, these include: the longitudinal modulus, the bulk modulus

(reciprocal compressibility), and Poisson's ratio. While the results of the

present study confirm this disparity, resonance-derived values of C^, B and

v are not reported here because of the large errors propagated in computing

them from E and G data. The basic reason for this disparity between dilatational-type and shear-type elastic constants is unclear.

However, a more fundamental problem is also unresolved -- how can the

anomalous elastic behavior of invar be understood, independent of the

resonance/pulse-echo disparity? Despite many attempts, no consensus of

thinking on the theory of invar exists. The AE^ effect due to linear magnetostriction due to domain walls accounts only for anomalies in shear-type

elastic constants. The AE effect due to volume magnetostriction accounts

oo only for anomalies in dilatation-type elastic constants. As described above,

both types of elastic constants are anomalous in invar. Thus, a combined

AE, - AE effect would have to be invoked. Hausch [5] disputed this approach

A oo and proposed that the Heisenberg magnetic exchange energy must be considered

for invar. Before these questions can be resolved, additional low-temperature

elastic measurements must be made on invar-type materials to permit quantitative

checks on proposed theories.

87
REFERENCES

1. W. S. McCain and R. E. Maringer, DMIC Memorandum 207 (1965).

2. S. J. Rosenberg, NBS Monograph 106, U.S. Govt. Print. Off. (1968).

3. G. Hansen and H. Warlimont, Z. Metallk. 64:152 (1973).

4. W. F. Schlosser, J. Phys. Chem. Solids 32:939 (1971).

5. G. Hausch, Phys. Status Sol idi (a) 15:501 (1973).

6. Handbook on Materials for Superconducting Machinery, Metals and Ceramics

Information Center, Battelle Columbus Laboratories, Columbus, Ohio (1974)

7. T. F. Durham, R. M. McClintock, R. P. Reed, C. J. Guntner, and K. A.

Warren, Cryogenic Materials Data Handbook, U.S. Dept. of Commerce (1961).

8. P. P. M. Meincke and J. Litva, Phys. Lett. 29A:390 (1969).

9. R. Fletcher, J. Phys. C 2:2107 (1969).

10. T. Maeda, J. Phys. Soc. Japan 30:375 (1971).

11. G. Hausch and H. Warlimont, Acta Met. 21 : 401 (1973).

12. G. Hausch and H. Warlimont, Z. Metallk. 63:547 (1972).

13. P. Chevenard, Trav. Mem. Bur. Int. Poids Mes . 27:142 (1927).

14. W. Koster, Z. Metallk. 35:194 (1943).

15. M. E. Fine and W. C. Ellis, Trans. AIME 188:1120 (1950).

16. W. H. Hill, K. D. Shimmin, and B. A. Wilcox, Amer. Soc. Test. Mater.

Proc. 61:890 (1961).

17. Y. Tino and T. Maeda, J. Phys. Soc. Japan 18:955 (1963).

18. H. M. Ledbetter and R. P. Reed, J. Phys. Chem. Ref. Data 2:531 (1973).

19. E. R. Naimon, W. F. Weston, and H. M. Ledbetter, Cryogenics 14:246 (1974)

20. W. F. Weston, J . Appl . Phys

.

, forthcoming.

21. R. M. Bozorth, W. P. Mason, and H. J. McSkimin, Bell System Tech. J.

30:970 (1951).

88
LIST OF TABLES

I. Summary of low-temperature elastic studies on invar.

II. Elastic constants of invar at selected temperatures.

LIST OF FIGURES

1. Temperature variation of the shear modulus of invar.

2. Temperature variation of the Young's modulus of invar.

3. Temperature variation of the longitudinal modulus of invar.

4. Temperature variation of the bulk modulus of invar.

5. Temperature variation of Poisson's ratio of invar.

89
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91
0.590

0.580

0.570

0.560

<= 0.540 —

§ 0.530 —

0.520 —

0.510

0.500

0.490 —

0.480

A Present resonance results O Present pulse-echo results

Maeda (H=1.9 kOe), right scale

Durham et al

.

150 200

TEMPERATURE (K)

Figure 1. Temperature variation of the shear modulus of invar.

— 0.460 — 0.450

350

150 200

TEMPERATURE (K)

Figure 2. Temperature variation of the Young's modulus of invar.

350

O

uoCO

I—

t

o

a.

150 200 250

TEMPERATURE (K)

300

Figure 5. Temperature variation of Poisson's ratio of inva;
Materials Science and Engineering, 20 (1975) 185—194

© Elsevier Sequoia S.A., Lausanne — Printed in The Netherlands

Dynamic Low-Temperature Elastic Properties of Two Austenitic Nickel—Chromium—Iron Alloys*

W.F. WESTON1", H.M. LEDBETTER and E.R. NAIMONt*

Cryogenics Division, Institute for Basic Standards, National Bureau of Standards, Boulder, Colorado 80302 (U.S.A.)

(Received in revised form March 3, 1975)

SUMMARY

The zero-magnetic-field low-temperature

elastic properties of two polycrystalline nickel - chromium - iron alloys were determined ultra- sonically between 4 and 300 K. Results are given for: longitudinal and transverse sound

velocities, Young's modulus, shear modulus,

bulk modulus, Poisson's ratio and elastic Debye

temperature. Effects of alloying are discussed.

The elastic property changes due to additions

of chromium and iron to nickel are reviewed

comprehensively.

INTRODUCTION

Elastic properties of metals at low temperatures have twofold interest. First, such information is essential for understanding the basic

aspects of mechanical deformation, which are often described by dislocation models that

contain the elastic constants explicitly. Second,

low-temperature elastic properties are essential

design parameters for cryogenic structures,

for predicting deflections due to any combination of stress and temperature.

Austenitic nickel-base alloys containing

chromium and iron are standard engineering

materials. They have high heat resistance, high

corrosion resistance, good high-temperature

strength, and can be readily fabricated into

structures. Many of these materials also have

low-temperature applications. In particular,

* Contribution of NBS, not subject to copyright. t NRC-NBS Postdoctoral Research Associate, 1973-4.

* Present address: Dow Chemical USA, Golden,

Colorado 80401.

because of its mechanical properties, Inconel*

600 is used extensively at cryogenic temperatures. Its strength increases with decreasing

temperature, while its ductility and toughness

are preserved. Inconel X-750 (formerly Inconel

X) is similar to Inconel 600, but it contains

small amounts of aluminum, titanium and

niobium. Aluminum and titanium make pre- cipitation hardening possible by forming

Ni3(Al, Ti) by suitable thermal treatment.

(Inconel 600 is usually hardened by low-temperature mechanical deformation.) Niobium

further stiffens the matrix through solidsolution hardening and stabilizes the carbides.

The zero-magnetic-field dynamic low-temperature elastic properties of Inconel 600 and

Inconel X-750 are reported here. A pulsesuperposition method was used to determine

ultrasonic wave velocities in specimens pre- pared from as-received commercial bar stock.

EXPERIMENTAL

Inconel 600 and X-750 alloys were obtained

from commercial sources in the form of about

3-in. (7.6-cm) and 4-in. (10.2-cm) diameter

rods, respectively. Cylindrical specimens

| -in. (1-cm) long and |-in. (1.6-cm) diam. were prepared by grinding. Opposite faces were flat and parallel within 100 X 10-6

in. (2.5 um). Chemical compositions (obtained

from mill analyses), hardness numbers and

mass densities are given in Table 1 . Hardnesses were determined by a standard method, and

mass densities were determined by hydrostatic

* Tradenames are used to characterize materials; they

are not NBS endorsements of particular products.

95

nHBHMSSQHHBHQSBHBBBRiSlli
186

TABLE 1 Compositions and properties of alloys

Alloy Chemical composition, mill analyses (wt.%)

Ni Cr Mn Fe S Si Cu C Al Ti Nb+Ta

Inconel Bal 60015.8 0.20 7.20 0.007 0.20 0.10 0.04 - -

Inconel 73.52

X-750

15.2 0.20 6.58 0.007 0.25 0.04 0.04 0.78 2.51 0.89

Alloy Hardness

(DPH No., 1 kg load) Mass density at 294 K

(g/cm3

)

Condition

Inconel 600 179 8.415 As received; hot rolled annealed

Inconel X-750 330 8.238 As received; hot rolled 1625°F

(1158 K) and aged

weighing using distilled water as a standard.

A pulse-superposition method [1] was used

to determine all ultrasonic velocities except

the shear mode in Inconel X-750. This particular

mode was highly attenuated, and for its mea- . surement a standard pulse - echo technique was used. All velocities were measured along

the rolling direction of the rod.

Quartz transducers (10 MHz) were bonded

to the specimens with phenyl salicylate for

room-temperature measurements and with a stopcock grease for lower temperatures. When

these bonds failed occasionally at very low

temperatures, a silicone fluid (viscosity = 200,000 cP at 25°C) was used for bonding.

The specimen holder is shown schematically

in Fig. 1. The holder was placed in the ullage

of a helium dewar and lowered stepwise to achieve cooling. Temperatures were monitored

with a chromel - constantan thermocouple

contacting the specimen.

Velocity measurements were made over the

range of 4 - 300 K for Inconel 600. For Inconel

X-750, however, measurements were made

over the range of 40 - 300 K. Below 40 K, the

ultrasonic attenuation for both longitudinal

and transverse waves was too high to permit

accurate velocity determinations to be made.

However, if a material is well behaved, little change is expected in its sound velocities from

40 to 0 K, and semi-theoretical curves fitted to

.

the data points (as explained in the next sec- tion) should give accurate low-temperature

values.

The experimental quantity of interest is the

transit time t for an ultrasonic wave to propagate from one end of the specimen to the other

and back. The ultrasonic wave velocity v is then given by

v = 2l/t, (1)

where / is the specimen length. Elastic moduli

C are related to ultrasonic velocities by

C = pv\ (2)

T hermoc oup 1

e

akelite cylinder containing

externally tunable impedancematching circuit

Electrical lead

akelite, guide

rass plunger, spring loaded Transducer

Spec i m e n

Brass base

Heat vane

Fig. 1. Specimen holder.
187

where p is the mass density.

Denoting room-temperature values by the

subscript r, it follows that

P/Pr = (/r/0

3

. (3)

Then eqn. (2) can be rewritten as C = Cr(/r //)Ur /f) 2

. (4)

Thus, the temperature dependence of an elastic modulus is determined from the change in transit time and from the length change due to thermal expansion or contraction. However,

the relative length change for Inconel over the

temperature range from 300 K to 4 K is only

about 0.23% [2]. These small thermal-contraction corrections are neglected here; their

omission introduces a maximum error also of

0.23% in the low-temperature values of velocity

and modulus. Such small errors are insignificant

for the purposes of the present study. Based on previous work in our laboratory, maximum

uncertainties in the velocity measurements are estimated to be about 1%.

A quantity of some interest is the temperature coefficient (l/C)(dC/dT), which is a measure of the relative modulus change due

to temperature changes. From eqn. (4), it follows that

(l/C)(dC/dT) = -2(l/f)(df/dT) - a (5)

where a = (l//)(d//dT) is the coefficient of

linear thermal expansion. For Inconel, a is 0.13 X 10~4 K

_1

at room temperature [2].

2.930 -

2.770 -

2 ;: l 1 1 i 1

0 50 100 150 200 250 300

TEMPERATURE (K)

Fig. 2. Temperature dependence of longitudinal modulus

Ccj = PVq of two nickel - chromium - iron alloys.

where C, s, and t are adjustable parameters and

T is temperature. The value of C at T = 0 K is C°, and —s/f is the high-temperature limit of

dC/dT. By invoking an Einstein oscillator

model of solids, it can be shown (in the absence

of electronic effects) that t is the Einstein

characteristic temperature. Parameters C°, s and t for Inconels 600 and X-750 are given in Table 2. Room-temperature values of the temperature coefficients of the elastic moduli are given in Table 3. Curves in Figs. 2 and 3 are plots of eqn. (8)

determined by an unweighted least-squares fit of the data. For Inconel 600, average percentage

differences between measured and curve values

RESULTS

Temperature dependences of the longitudinal

and transverse moduli are shown in Figs. 2 and

3. These moduli are given by

Ci = pv% = B+±G, (6)

and

Ct = pv

2

t = G. (7)

Here yfi and vt are the longitudinal and trans- verse sound-wave velocities, p is the mass

density, B is the bulk modulus, and G is the

shear modulus.

Temperature dependences of both Ce and Ct were least-squares fitted to a theoretical re- lationship suggested by Varshni [3]:

C = C°-sl(etlT -l), (8)

0.362

0.B52 -

0 50 100 150 200 250 TEMPERATURE (K)

Fig. 3. Temperature dependence of transverse or shear modulus Ct = puf = G of two nickel - chromium - iron alloys.

97
188

TABLE 2 Parameters in the temperature-dependence eqn. (8)

Alloy Mode C°

(10 11 N/m2

)

s(10 uN/m2

)

t(K)

Inconel pvl 2.909 0.107 203.1 600

P«? 0.828 0.063 246.0

Inconel 2.920 0.105 165.9

X-750

2 P"t 0.846 0.120 315.8

are 0.05% for both the longitudinal and trans- verse moduli. For Inconel X-750, the average

percentage differences between measured and

curve values are 0.04% and 0.18% for the

longitudinal and transverse moduli, respectively.

The comparatively large error for the transverse modulus of Inconel X-750 was due to a rela- tively poor echo pattern resulting from ultra- sonic-wave scattering. There are many possible sources of this scattering — for example lattice inhomogeneities and crystalline anisotropy.

While polycrystalline aggregates (quasi-isotropic solids) have only two independent elastic constants, several constants are commonly

used for various applications. The four most

common are the bulk modulus B, Young's

modulus E, the shear modulus G and Poisson's

ratio v. The relationships among these are: 1 =

1

- +

JL

- (9) E 3G 9B 1 '

and

P=fc

^-1. (10)

These elastic constants were calculated from

the moduli shown in Figs. 2 and 3 by the

2.202 -

2.026 -

0 50 100 150 200 250 ,300 TEMPERATURE ( K)

Fig. 4. Temperature dependence of Young's modulus

of two nickel - chromium - iron alloys.

relationships:

£ = 3Ct (Cv -fCt )/(Cc -Ct ), (11)

5 = C, -|Ct , (12)

and

i/ = !(Ce -2Ct )/(Ce -Ct ) (13)

where Cc and Ct are the least-squares values

obtained from fitting the experimental data

with eqn. (8). The constants E, B and v are shown in Figs. 4 - 6.

It is of interest to calculate the elastic Debye

temperature 6 for the two alloys. This fundamental parameter is important in the lattice properties of solids and is related to the

average elastic wave velocity by [7]

0 = K(u), (14)

TABLE 3 Temperature derivatives of elastic constants at room temperature (10—

4

K—1

)

Source Material 1 <±B 1 d£ 1 dG 1 dv

Bdf Edf Gdf udT

Present Inconel 600 —1.20 —2.95 —3.21 1.10

Present Inconel X-750 —1.02 —4.07 —4.53 1.97

Ref. 4 Nickel* —1.56 -4.16 —4.53 1.72

Ref. 5 Inconel X —3.09

Ref. 6 Inconel X -2.67

Ref. 12 Inconel X -3.16

* Based on single-crystal data averaged by the Voigt - Reuss - Hill method.

98
189

1.820

1.310 -

1.730 -

1.720 ' 1 1

1 1 1 1 0 50 100 150 200 250 300 TEMPERATURE ( K)

Fig. 5. Temperature dependence of bulk modulus

(reciprocal compressibility) of two nickel - chromium - iron alloys.

where

Here h is Planck's constant, k is Boltzmann's

constant, N is Avogadro's constant, p is the mass density, and A is the effective atomic

weight. The average velocity is given by

m^2^y^ (16)

The Debye temperature of each alloy at T = 0 K,

and also of nickel, is given in Table 4.

t'-3i2i 1 r

0.310 -

0.309 -

0.26; I 1 1 : 1 1 0 50 inn 150 200 250 300 TEMPERATURE (K)

Fig. 6. Temperature dependence of Poisson's ratio of two nickel - chromium - iron alloys.

TABLE 4

Elastic Debye temperatures at T = 0 K

Source Material 0(K)

Present Inconel 600 464.9

Present Inconel X-750 473.5

Ref. 8 Nickel 476.0

DISCUSSION

The elastic properties of both Inconel 600

and Inconel X-750 behave regularly with respect to temperature. The elastic moduli (Cg, Ct = G,

B, E) decrease with increasing temperature,

show a relative flatness at low temperatures,

achieve zero slope at T = 0 K, and approach

linear behavior at high temperatures. Poisson's

ratio also behaves regularly, having a positive

temperature coefficient. For both the bulk

modulus and Poisson's ratio, Inconel X-750

shows unusual curvatures in the modulus - temperature curves. The reason for this is not

understood; atomic ordering based on the

Ni3 (Cr, Fe) composition is a possible explanation.

Ordering often occurs on cooling and usually causes an increase in the elastic stiffnesses. Magnetic ordering might also affect the elastic stiffnesses. Assuming that the specimens studied are representative of the two alloys, then conclusions concerning their relative elastic behavior can be drawn. Not surprisingly, as shown in Figs. 2 - 6 and in Tables 2 and 3, for

most practical purposes the two alloys are

elastically identical. Thus, additions of small

amounts of aluminum, titanium and niobium

have little effect on the elastic properties of

nickel alloys containing about 15% chromium

and about 7% iron.

Very few elastic data exist for these alloys

[5,6,9 - 13], especially Inconel 600. Most in- formation has appeared in engineering reports

and is summarized in refs. 9-13. For comparison the room-temperature values of E, G,

B and v are given in Table 5. Good agreement

is observed between previous and present

results. The polycrystalline elastic constants

of nickel are also given in Table 5, both those

determined from experiments on polycrystalline

aggregates [4] and those obtained from singlecrystal data [15,16] using a Voigt - Reuss - Hill average. The elastic Debye temperatures of

99
190

TABLE 5 Room-temperature elastic constants of Inconel 600 and Inconel X-750; units of 10 1:l N/m2 except

v (dimensionless)

Source Inconel 600 Inconel X-750

E G B v E G B V

Present 2.036 0.780 1.760 0.307 2.040 0.784 1.736 0.304

Ref. 5 0.789

Ref. 6 2.144

Refs. 9,10 2.11 0.75 2.11 0.75 0.29

Ref. 11 2.075 0.714 0.29

Ref. 12 2.041 0.824

Ref. 13 2.109 2.109 0.749

Nickel - 26 iron

Ref. 14 2.059 0.792 1.731 0.302

Nickel (polycrystal, zero magnetic field)

Ref. 4 1.97 0.785 1.903 0.296

Nickel ( VRH-averaged single crystal, zero magnetic field)

Ref. 4 2.10 0.799 1.863 0.312

/

the alloys are also within 2% of the value cal- culated from the single-crystal elastic data of

nickel.

Low-temperature elastic data have been

reported only once previously for either of

these alloys; based on static measurements,

Young's modulus and the shear modulus of

Inconel X-750 were reported between room

temperature and liquid-hydrogen temperature

(20 K) [12]. The reported Young's modulus

agrees closely with present results while the

reported shear modulus is about 4% higher

than the present results and has a lower roomtemperature temperature derivative.

It is emphasized that the data reported here

are dynamic (adiabatic) rather than static (isothermal) and apply to rapid, rather than

slow, loading. In most cases the differences

between adiabatic and isothermal elastic

constants are small. The conversion formulas

can be written as [17]

ES = ET (1 + ET Ta 2 /Cp p), (17)

BS = BT (1 + 9BT Ta 2 /Cp p), (18)

vs = v T + (1 + vT )ET Ta 2ICp p, (19)

and

GS = GT , (20)

where subscripts S and T denote adiabatic and

isothermal, respectively. In the above equations

T is the absolute temperature, a is the linear

thermal expansion coefficient, Cp

is the heat

capacity per unit mass at constant pressure, and p is the mass density. For the Inconel

alloys at room temperature it is found that

(Es - ET )/ES = 0.003, (BS -BT)/BS = 0.022,

(t>s - vT)/vs = 0.012 and (Gs — GT ) = 0. (21)

At lower temperatures the differences are smaller; they vanish at T = 0 K.

Alloying effects on material properties can

be analyzed many ways. The simplest approach

is a linear mixture model:

C= E x^, (22)

i= 1,N

where, in present context, C represents a general elastic constant of the alloy, which has N

components of species i; x; are the weight

fractions, and Cj are the elastic constants of

the unalloyed species. This model works best

for mixtures of similar metals where effects of

atomic ordering, phase transformations, etc. are absent. It can be shown simply that this model fails for transition-metal alloys, even for

simpler, binary systems such as iron - nickel.

A better model proceeds on the basis of

100
191

alloy data rather than component data:

£ (I^W (23)

where C0is the elastic constant of the reference

material, usually (but not necessarily) a pure

component, (3C/3Xj) are slopes of the C - curves, and are the weight fractions. Expressions like eqn. (23) have been used pre- viously to describe the elastic properties of

ternary alloys [18]; it assumes that alloying

effects (AC/Ax) are composition independent

in the range of interest and are linearly additive.

Considering Inconel alloys to be, in a first approximation, alloys of iron and chromium in nickel, then three data are required (besides

composition) to solve eqn. (23) — the elastic constants of nickel and the changes of these

constants with iron and with chromium additions. Nickel's elastic constants are well

known [4].

Effects of iron additions on the elastic properties of nickel have been studied extensively [19 - 32], although the subject has ap- parently not been reviewed. A summary of the

results is given in Table 6, which gives the

logarithmic composition derivatives (1/C)(AC/

Ax), where C represents an elastic constant

such as E, G, B or v, and x denotes weight percent iron alloyed into nickel. The effects of

alloying iron into nickel are best defined for

the Young's modulus and the bulk modulus; the

change of the bulk modulus is quite large.

Effects of chromium additions on the elastic properties of nickel have been reported only

by Chevenard [21] and are shown in Table 7.

If the value of (l/G)(AG/Ax)Fe from Table 6

is substituted into eqn. (23) along with the

other input data, a value of (l/G)(AG/Ax)Cr = —6.6 is calculated. Since this disagrees with

Chevenard's value of 2.9, the model represented

by eqn. (23) also fails for the alloys of interest

here.

The necessity of invoking higher-order

terms in the Taylor expansion is indicated:

+ i £ (|£)a*,

O 0 <-o i= 1 ,N \ oxi/

+ M&Ax^' + - (24)

In particular, the interaction term in this case:

c\{d^Jh) XciXFe (25)

may be quite important since alloying alters 3d-shell electronic structure, which in turn

affects cohesion and elastic properties. At

TABLE 6

Effects of iron on the elastic properties of nickel at room temperature (evaluated from 0 - 7% iron) —

1 (AC\ —- 3 xlO d C \Ax

E G B V

Honda (1919) 2.4 1.4 Honda, Tanaka (1926) 0.3 0.7 1.7 Chevenard (1927) 0.0 Nishiyama (1929) 4.1 Ebert, Kussmann (1937) -15.0

Forster, Koster (1937) 1.7 Engler (1938) 1.0 Koster (1940) 1.9 Chevenard, Crussard (1943) 0.0 Koster (1943) 1.7 Yamamoto (1959) 7.8 Sakurai et al. (1964)* —0.4

Shirakawa et al. (1969)* 5.3 9.8 —17.5 -17.3

Tanji et al. (1970) 0.8 2.3 -16.8 -10.7

Average observed values 2.4 2.3 —16.4 -8.8

Linear mixture model -0.6 -0.5 —1.1 -0.3

* Single-crystal data were averaged by the method of Voigt.

101
192

TABLE 7 Effects of chromium on the elastic properties of nickel

Source 1 (ae\ WAG\

G\Axf\xlO- B\Axf xl(T v\Ax 1x10°

Linear mixture

model

Chevenard (1927)

2.2 3.1

2.9

-1.1 -3.2

present, data are insufficient for evaluating

this interaction term. Chevenard 's [21] data on Ni - Cr - Fe alloys indicate that for the

shear modulus the interaction term might be

as large as 40 X 10~3x FexCt , which is large

relative to the first-order coefficients in Tables

6 and 7. It is difficult to make any quantitative

estimate of the second-order mutual interaction terms (Fe - Fe, Cr - Cr). But from the

existing data they seem to be smaller than the

cross-term (Fe - Cr). Additional evidence for a

large interaction term results from the following

reasoning. Since the elastic properties of the

two Inconel alloys are so similar, additional

alloying elements beyond chromium and iron seem to have negligible effects or to cancel. In either case, considering Inconels to be ternary

alloys of nickel, chromium and iron seems

justified from the viewpoint of their elastic properties. However, both alloys behave elas- tically very much like a nickel - iron alloy

with the same nickel content; chromium behaves as if it were iron. This is shown in Table 5, which includes the elastic properties

of a nickel-26 iron alloy, which were obtained

from single-crystal data by using a Voigt - Reuss - Hill-arithmetic average. Why chromium should

behave like iron is not understood.

Another anomaly of the Ni - Cr - Fe system

is that the elastic stiffnesses of nickel are in- creased by either chromium or iron additions.

Usually the elastic stiffness of a host metal is reduced by alloying. The reverse effect in this case could be magnetic in origin, as discussed

below.

Magnetic interactions contribute to the

energy of a material and therefore also to its elastic constants, which are related to the

second spatial derivative of the total energy.

However, extra energies due to magnetism

take many forms, and several models have been

suggested to account for the anomalous elastic

behaviour of magnetic materials. At least two

of these models are relevant here.

Khomenko and Tseytlin [33] showed for

alloys in the invar region (Fe ~ 35 Ni) that

chromium additions increased Young's modulus

with (1/E)(AE/Ax) having a value near 30 X 10~s

, a very large effect compared with

the coefficients in Tables 6 and 7. These authors

showed that chromium increases E because it suppresses the usual AE effect AEX , owing to

linear magnetostriction, which reduces E. Since

AEx is proportional to a magnetostriction

constant X that is quite large for nickel, this

is a possible model for the anomalous elastic behaviour of nickel - chromium - iron alloys.

Hausch [34] used a molecular-field approximation to evaluate the Heisenberg exchangeenergy J(r) contribution to the elastic constants

of magnetic materials. For nickel it was shown

that d

2J(r)/dr2

is negative (therefore the elastic constant contribution is negative), while for

face-centered cubic iron the derivative is positive. The derivative has not been evaluated

for chromium; it is probably positive as deduced

from its approximate position on a Bethe - Slater curve. Thus, it is clear qualitatively that

Hausch 's model might also be able to explain

the anomalous elastic behaviour of these alloys.

The two models could be tested by lowfrequency elastic measurements on appropriate

alloys, with and without a saturating magnetic

field. The AEX variation with alloying disappears in a saturating magnetic field since the

magnetic domain walls are immobilized.

Finally, a speculative suggestion is made

concerning the source of the very high acoustic

attenuation in Inconel X-750 below 40 K. Iron

and nickel have Curie temperatures of 1040 K

and 630 K, respectively. Alloying with

chromium reduces both of these drastically;

about 50% chromium in iron suppresses completely the paramagnetic - ferromagnetic

102
193

transition; for nickel, fewer data exist, but only

about 12% chromium may be required to suppress the Curie transition in nickel. No

studies seem to have been reported on the

Curie temperatures of Ni - Cr - Fe alloys [35].

Thus, a paramagnetic - ferromagnetic transition

in Inconel X-750 at low temperatures is possible. If it occurs, the exact transition temperature should depend strongly on chromium

content. Such a transition would explain the

attenuation increase as due to acoustic-wave

scattering from ferromagnetic domain walls. The attenuation would not be associated with

the Curie transition itself, but with a change

in the magnetic state, a state produced by the

Curie transition. Simple dip tests using a small

permanent magnet indicated that both Inconels

600 and X-750 are ferromagnetic at liquidnitrogen temperatures. Tests for ferromagnetism between room and nitrogen temperatures were not made. Further studies on this problem are planned. Studies that would be valuable

include low-temperature acoustic wave-velocity

and attenuation measurements in a saturating

magnetic field. If a ferromagnetic transition occurs, then lower-frequency elastic measure- ments would be useful for determining the

so-called AE effect. No AE effect is seen in

the MHz region because ferromagnetic domain

walls cannot respond to a high-frequency

mechanical stress [36]. Measurements of Curie

temperatures and saturation magnetic moments

would also be useful.

CONCLUSIONS

From the results of this study the following

conclusions are drawn:

1. The elastic properties of Inconel 600 and

Inconel X-750 are quite similar.

2. At lower temperatures, Inconel X-750

has a slightly higher shear modulus and Young's

modulus than has Inconel 600, but a slightly

lower bulk modulus.

3. At low temperatures, Poisson's ratio is slightly lower for Inconel X-750 than for

Inconel 600.

4. Inconel 600 shows a regular temperature

dependence in all its elastic properties.

5. The temperature dependence of the elastic properties of Inconel X-750 is regular except

for the bulk modulus, which has a region of

positive curvature, and Poisson's ratio, which

has a minimum near 80 K.

6. Inconel X-750 highly attenuates 10 MHz

sound waves, both longitudinal and transverse,

below about 40 K. This is speculated to result

from a change in the magnetic state, which

results from a paramagnetic-to-ferromagnetic

transition that occurs above 77 K. Why Inconel

600, which also undergoes a magnetic transi- tion, has lower attenuation is not understood.

7. The composition dependence of the

elastic properties of these alloys cannot be described by a simple model, probably because

of d-shell and magnetic interactions.

ACKNOWLEDGEMENTS

This work was supported in part by the

Advanced Research Projects Agency of the

U.S. Department of Defense. Dr. R.L. Moment

of Dow Chemical contributed a critical reading

of the manuscript.

REFERENCES

1 H.J. McSkimin, J. Acoust. Soc. Am., 33 (1961) 12.

2 R.J. Corruccini and J.J. Gniewek, Thermal Expansion of Technical Solids at Low Temperatures,

NBS Monograph 29, 1961, p. 10.

3 Y.P. Varshni, Phys. Rev., B2 (1970) 3952.

4 H.M. Ledbetter and R.P. Reed, J. Phys. Chem.

Ref. Data, 2 (1973) 531.

5 H.C. Burnett, unpublished results, 1956.

6 W.H. Hill, K.O. Shimmin and B.A. Wilcox, Am.

Soc. Testing Mater. Proc, 61 (1961) 890.

7 P. Debye, Ann. Phys. (Leipzig), 39 (1912) 789.

8 R. Wanner, Can. J. Phys., 48 (1970) 1270.

9 Intern. Nickel Co., Rept. on Inconel Alloy 600,

1973.

10 Intern. Nickel Co., Rept. on Inconel Alloy X-750,

1970.

11 NERVA Program Materials Properties Data Book,

Vol. 1A, Nickel Base Alloys, Aerojet Nuclear

Systems, Sacramento, Calif., 1970. 12 Cryogenic Materials Data Handbook, AFML Rept. No. ML-TDR-64-280, PB 171809, revised, Aug.

1964. 13 Aerospace Structural Metals Handbook, Vol. II, AFML. Rept. No* ASD-TDR-63-741, Mar. 1966. 14 N.G. Einspruch and L.T. Claiborne, J. Appl. Phys.,

35 (1964) 175. 15 S.G. Epstein and O.N. Carlson, Acta Met., 13 (1965) 487. 16 G. Simmons and H. Wang, Single Crystal Elastic Constants and Calculated Aggregate Properties:

A Handbook, M.I.T. Press, Cambridge, 1971, p. 223. 17 L.D. Landau and E.M. Lifshitz, Theory of Elasticity, Pergamon Press, London, 1959, p. 17. 18 J.B. Greer and E.H. Bucknall, Trans. Am. Soc.

Metals, 57 (1964) 554.

103
194

19 K. Honda, Sci. Rept. Tohoku Univ., 8 (1919) 59. 20 K. Honda and T. Tanaka, Sci. Rept. Tohoku Univ.,

15 (1926) 1. 21 P. Chevenard, Trav. Mem. Bur. Int. Poids Meas., 27 (1927) 116. 22 Z. Nishiyama, Sci. Rept. Tohoku Univ., 18 (1929)

539. 23 H. Ebert and A. Kussmann, Physik. Z., 38 (1937)

437. 24 F. Fdrster and W. Kdster, Naturwiss., 25 (1937) 436. 25 O. Engler, Ann. Physik, 31 (1938) 145. 26 W. Kdster, Arch. Eisenhuettenw., 6 (1949) 271. 27 P. Chevenard and Ch. Crussard, Compt. Rend.,

216 (1943) 685. 28 W. Kdster, Z. Metallk., 10 (1943) 194. 29 M. Yamamoto, Sci. Rept. Tohoku Univ., 11 (1959) 102.

30 J. Sakurai, M. Fujii, Y. Nakamura and H. Takaki,

J. Phys. Soc. Japan, 19 (1964) 308. 31 Y. Shirakawa, Y. Tanji, H. Moriya and I. Oguma,

J. Japan Inst. Metals (Sendai), 33 (1969) 1196. 32 Y. Tanji, Y. Shirakawa and H. Moriya, Rept. No.

1462, Res. Inst, for Iron, Steel and Other Metals

(1970), p. 84; also in J. Japan Inst. Metals (Sendai),

34 (1970) 417. 33 O.A. Khomenko and A.M. Tseytlin, Fiz. Metal,

i Metalloved., 28 (1969) 246. 34 G. Hausch, Phys. Status Solidi, 15 (1973) 501. 35 T.F. Connolly (ed.), Bibliography of Magnetic

Materials and Tabulation of Magnetic Transition Temperatures, Vol. 5, Solid State Physics Literature Guides, Plenum Press, New York, 1972.

36 W.P. Mason, Phys. Rev., 83 (1951) 683.

104
Low-temperature elastic properties of four austenitic

stainless steels

H. M. Ledbetter, W. F. Weston,n and E. R. Naimon*+

Cryogenics Division, Institute for Basic Standards, National Bureau of Standards, Boulder, Colorado 80302

(Received 11 February 1975)

The elastic properties of four austenitic stainless steels—AISI 304, AISI 310, AISI 316, and A286—are reported over the temperature range 300-4 K. These properties include longitudinal modulus, shear modulus, Young's modulus, bulk modulus (reciprocal compressibility), Poisson's ratio, and elastic Debye

temperature. Elastic constants were determined from measurements of longitudinal and transverse sound- wave velocities using an ultrasonic (10 MHz) pulse-superposition method. Measurements were made in the absence of a magnetic field; these alloys undergo paramagnetic-to-antiferromagnetic transitions at low

temperatures. For all four alloys, the shear modulus behaves regularly with respect to temperature. The

other elastic constants, all of which have a dilatationa! component, decrease anomalously at temperatures below 80 K. The largest anomaly, about 3%, is in the bulk mudulus of the 304 alloy; this modulus is lower at 0 K than at 300 K. Results are interpreted on the basis of the Doring effect, which results from a

large volume magnetostriction in the magnetic phase. This may be the first report of a Doring effect in antiferromagnetic materials.

PACS numbers: 62.20.D, 65.70., 75.80.

I. INTRODUCTION

Austenitic stainless steels are attractive materials

for mechanical applications at low temperatures. Primarily, this is due to their having, at room temperature, a face -centered cubic (fee) crystal structure. Metals having this crystal structure usually do not be- come brittle at lower temperatures. In general, steels

that remain austenitic at cryogenic temperatures show

increased tensile strength, a smaller increase in yield

strength, and little change in ductility . However, it is not axiomatic that fee materials will always perform

well at lew temperatures. For example, lower temperatures may promote a change of the crystal structure from fee to body -centered cubic (bec) or to closepacked hexagonal (cph), thus probably embrittling the material. Other changes such as atomic ordering or magnetic ordering may also occur at low temperatures;

these also effect mechanical behavior. Thus, the nature

and the magnitude of a material's low-temperature

properties cannot be predicted a priori from roomtemperature observations, and there is no substitute for careful low -temperature experimental determinations of the important properties of each material of interest. Gilman1 concluded that "the most important

mechanical characteristic of a crystal is its elastic modulus".

The same elastic constants that are related to fundamental interatomic forces in solids are also used in engineering design. For example, Poisson's ratio is an essential design parameter in problems of plate

buckling or of pressure -vessel design. Young's modulus and Poisson's ratio are required if plane -stress

data and plane -strain data are to be interconverted, a

technique used often in the elastic stress -strain analysis of solids.

In this paper, the dynamic zero-magnetic -field elastic properties of four austenitic stainless steels —commonly designated AISI 304, AISI 310, AISI 316, and

A286 2 —are reported between 300 and 4 K. These properties include longitudinal modulus, shear modulus,

Young's modulus, bulk modulus (reciprocal compressibility), and Poisson's ratio. These elastic constants were determined dynamically, by measuring the velocity of longitudinally polarized and transversely polarized

ultrasonic (10 MHz) pulses propagating through polycrystalline specimens of commercial as -received alloys. At low temperatures, elastic anomalies were observed in

all four materials. These are believed to be associated

with transitions to antiferromagnetic states A magnetoelastic interpretation of the anomalies is given.

II. EXPERIMENTAL

A. Specimens

Materials were obtained from commercial sources

in the form of ] -in. (1. 9-cm)-diam rods. Their chemical compositions are given in Table I. Hardness and mass -density data on the alloys are given in Table II. Hardnesses were measured by standard metallurgical

methods, and mass densities were measured by

Archimedes 's method using distilled water as a standard. Materials were tested in their as-received conditions. Samples were prepared by grinding cylinders

TABLE I. Chemical analyses of the alloys, wfi, obtained from mill analyses.

Alloy Al C Cr Cu Mo Mn Ni P S Si Ti V Fe

304 0. 02 18.4 1.4 9. 7 0. 02 0. 01 0.6 balance 310 ~~0. 08 24. 8 0.1 0.1 1.7 20.8 0. 02 0. 02 0.7 balance

316 0. 05 16.8 0.2 2.1 1.9 11.7 0. 03 0. 02 0.4 balance A286 0.2 0. 04 14. 8 1.2 1.4 25.4 0. 01 0. 01 -0.6 2.1 0.3 balance

3855 Journal of Applied Physics, Vol. 46, No. 9, September 1975 Copyright © 1975 American Institute of Physii

105

3855
TABLE II. Densities and hardnesses of the alloys.

Alloy Mass density

(g/cm3

) Hardness

(DPHN, 1 kg load)

304 7 ftft 99 c 310 7.85 220

316 7.97 210 A286 7.95 275

^ in. (1.2 cm) thick with faces flat and parallel within

10-4 in. (2.5 u m).

B. Procedures A pulse -superposition 3 method was used to measure

the longitudinal and transverse sound-wave velocities between room temperature and liquid-helium temperature. The specimen holder, which was described previously4

, was placed. in the ullage of a helium Dewar and

lowered (raised) stepwise to achieve cooling (heating).

Measurements were made semicontinuously on cooling,

and a few points were checked on heating to verify

reversibility. Temperatures were measured by a chromel-constantan thermocouple contacting the speci- men. Quartz transducers (10 MHz) were bonded to the specimens with phenyl salicylate for room -temperature

measurements and with stopcock grease for lower

temperatures. No bond corrections were made since these are insignificant for present purposes. No ther- mal-contraction corrections were made; for the alloys

of interest this introduces a maximum error of 0.3%

over the 300 K temperature range. Maximum uncertainties in the absolute velocity measurements are esti- mated to be about 1? - The imprecision in the relative velocities is a few parts in 105

.

III. RESULTS

Longitudinal and transverse moduli are shown in

Figs. 1 and 2 as a function of temperature, The longitudinal modulus C, is given by

C,= P

v

2v (1) where p is the mass density and v l is the longitudinal wave velocity. The transverse modulus C t is identically

equal to the shear modulus G and is given by

C t =G=p

v

2

t

, (2) where v t is the transverse wave velocity. Young's modulus E is given by

£ = 3C( (C,-4C t )/(C,-C ( ) (3) and is shown in Fig. 3. The bulk modulus B, or reciprocal compressibility, is given by

B = C,-$Ct

, (4) and is shown in Fig. 4. Poisson's ratio v is given by

v=j;(C,- 2Ct )/(C l -Ct ) (5)

and is shown in Fig. 5. Temperature dependences of both C; and C t were

fitted by a least-squares method to a theoretical relationship suggested by Varshni5

:

C = C°- s/(e t ' T -l), (6)

where C°, s, and / are adjustable parameters and T is temperature. The value of C at T = 0 K is C°, and -s/t

is the high -temperature limit of the temperature derivative dC/dT, By invoking an Einstein oscillator model of solids, it can be shown (in the absence of electronic

effects) that t is the Einstein characteristic temperature, Parameters C°, s, and t are given in Table in. The average differences between the Varshni -curve

values and measured values were 0,04 and 0,05% for C; and Ct respectively, Of course, the low-temperature elastic anomalies are not described by Eq. (6),

FIG. 1. Temperature dependence of the longitudinal moduli of four stainless-steel

alloys.

150 200 250

TEMPERATURE (K)

3856 J. Appl. Phys., Vol. 46, No. 9, September 1975 Ledbetter, Weston, and Naimon 3856

106
which describes the extrapolated higher-temperature behavior shown as dashed lines in the figures s Temperature coefficients of the elastic constants at room temperature are given in Table IV.

Elastic Debye temperatures at 7 = 0 K were calculat- ed for both the antiferromagnetic and the extrapolated

paramagnetic states, and these are given in Table V ; The elastic Debye temperature 9 is related to the average sound-wave velocity according to e=K(v), (7) where

K={h/k){3No/^AY'\ (8) Here h is Planck's constant, k is Boltzmann's constant^

A" is Avogadro's number, p is the mass density, and A

is the effective atomic weight. The average velocity is given by

(v ) = [±(v

-3 + 2v-t 3 )]-''\ (9)

For comparison, the elastic Debye temperatures at T = 0 K of iron, chromium, and nickel are also included

in Table V,

IV. DISCUSSION

The elastic constants of all four materials exhibit regular behavior from room temperature to about 80 K or

lower. Below about 80 K, anomalous changes in the

elastic constants occur for all four materials. These

anomalies occur only in the elastic constants that have

a dilatational component —the longitudinal modulus, the

2 450

2 378

2 306

Ez 2 234

~o 2 162

in3

-J- 209O Q

O

2018

in

> 946 O>-

1 874

1 802

1 730

^-^^^^-^286

^310 ~

50 150 200 250 TEMPERATURE (K) 300

1.880

1.834

1.788

1.742

1.696

1.650

1.604

1.558

1.512

1.466

I 420

111

^310

^-A286

'

1

ti__304^^^

1 ' 50 100 150 200 250 TEMPERATURE (K) 300

FIG. 3. Temperature dependence of the Young's moduli of four

stainless-steel allovs. FIG. 4. Temperature dependence of the bulk moduli (reciprocal

compressibilities) of four stainless-steel alloys.

J. Appi. Phys., Vol. 46, No. 9, September 1975 107

Ledbetter, Weston, and Naimon 3857
0 314-

0.266 -

0 26o' 1 ' 1 1 1 1 0 50 100 150 200 250 300 TEMPERATURE (K) FIG. 5. Temperature dependence of Poisson's ratios of four

stainless-steel alloys.

bulk modulus, Young's modulus, and Poisson's ratio. No anomalies occur for the shear modulus.

The elastic properties of some of these alloys have

been studied previously at low temperatures. For example, stress -strain tests for E and G at 295, 77, and

4 K indicated anomalies somewhere below 77 K for both AISI 302 and AISI 303, but not for AISI 310. 6 Resonance

tests on AISI 303 for E and G simultaneously at temperatures between 320 and 4 K showed anomalies below

80 K. 7 All these anomalous results can be interpreted in terms of the usual A£ = AEX effect (discussed below) and

are not directly related to the low -temperature elastic anomalies reported in the present work.

The most likely causes of the anomalies are magnetic

transitions. Stainless -steel -type alloys have been shown

to be antiferromagnetic at lower temperatures; Neel

temperatures near 40 K have been reported6 - 8 for 304-

type alloys. However, no Neel transition was detected

for a 310-type alloy cooled to liquid-helium temperature

'; This suggests that effects due to local, rather

than long-range, magnetic ordering may be occurring

in these alloys. This would also account for the differ- ence in the 304 alloy between the Neel temperature

(40 K) and the temperature below which anomalous elastic behavior is observed (80 K). Short-range magnetic

order effects have been discussed by Schlosser9 for fee iron-nickel alloys. Changes of crystal structure are

probably not the cause of the anomalies since the changes were observed to be reversible within experimental error; changes of elastic constants due to crystal-structure changes are generally irreversible, showTABLE III. Parameters in the temperature-dependence equation, Eq. ((!).

Alloy Mode L(in 11 N/m2

)

s

(10H N/m2

)

t(K) 304 / 2. 608 0.297 371 8

> / 0. 803 0.089 251 8 310 I 2. 647 0.175 271 8

t 0. 767 0.051 173 7 316 1 3, 028 0.332 369 7

t 0. 914 0. 099 243 5 A286 I 2 . 636 0.334 454 4

t 0. 809 0. 080 243 1

3858 J. Appl. Phys., Vol. 46, No. 9, September 1975

TABLE IV. Temperature-Coefficients of the elastic constants

at room temperature AO"4 K"1 ).

Alloy

1 dB

B dT

1 dE

E dT

1 dG G dT

1V dvdT

304 -1.69 -4.15 -4. 54 1 78 310 -1.39 -3.71 -4.05 1 45

316 - 1.41 -4.17 -4.61 . 1 89 A286 -1 .22 -3. 77 -4.18 1 81

ing a particularly large hysteresis in iron-base alloys.

The elastic constants reported here showed a reversible

behavior and indicate that the magnetic transition occurs smoothly over a range of temperatures rather than at a unique transition temperature. Thus, the transition seems to be of the second-order type. 10

The usual so-called A£ effect, better designated A£x , due to linear magnetostriction strains induced by an applied stress, and usually interpreted by invoking a Weiss domain model (with domain rotations and domainwall motions), can also be excluded as a possible source

of the anomalous elastic behavior. Linear magnetostriction affects Young's modulus, the shear modulus, and

Poisson's ratio, but it does not affect the bulk modulus

because no magnetomechanical process will respond to the application of a hydrostatic stress. 11 Also, effects due to A£x are not observed at high frequencies12 because the domain -wall displacements cannot follow the applied stress, and therefore cannot contribute an additional strain that lowers the observed elastic stiffness. Thus, for present purposes, a high-frequency applied stress is essentially equivalent to applying a saturating

magnetic field; both nullify domain-wall contributions to the strain.

In the presence of a saturating magnetic field (or its effective equivalent), the only magnetic effect that alters E, £, and v without affecting G is the effect due to spontaneous volume magnetostriction. This effect is designated A£u , where u = (V -Vfi )/ Vp is the spontaneous vol- ume magnetostriction, V is the total volume, and V

p is the volume in the paramagnetic state. The effect was

first reported by Engler13 for an Fe-42 Ni alloy. It was

explained first by Doring14 using a thermodynamic analysis. Herein, this phenomenon will be called the Doring

effect. Doring reasoned that in the paramagnetic region

the elastic constant is measured at constant magnetization M, while in the ferromagnetic or antiferromagnetic

TABLE V. Elastic Debye temperature at T-0 K.

Alloy e (k) magnetic

e (K) extrapolated

paramagnetic

304 467.8 468.1 310 459. 0 459. 0 316 496. 7 496.8 A286 467.6 467. 7 Iron (bee) a 472.4 Chromium b 453.0

Nickel 1

' 476.0

aR. Wanner, Can

b F.H. Herbstein,

J. Phys. 48, 1270 (1970). Adv. Phys. 10, 313 (1961).

Ledbetter, Weston, and Naimon 3858 108
region the elastic constant is measured at constant magnetic field H. Doring showed that the magnitude of the anomaly in the Young's modulus is given by

where dco/dH is the forced volume magnetostriction due

to the magnetization, dM/dH is the high-field susceptibility x, £

2 is EUE H and a is the tensile stress. An al- ternative derivation of Eq. (10) was given by Hausch. 15 [Wohlfarth16 stated, without proof, that the factor of §

in Eq. (10) should be replaced by (1 - 2v)2 where v is Poisson's ratio. For the present materials this intro- duces a factor of approximately 2, which is insignificant for present purposes. ] It should be noted that the Doring

effect always lowers the moduli {EM > EH), whether the volume magnetostriction is positive or negative. This is related to the fact that a relaxation mechanism is in- volved in going to an antiferromagnetic state, and such

mechanisms always soften the elastic stiffness. The

Dbring effect in iron-nickel alloys was discussed by

Koster, 17 who denoted it as A£x rather than as A£w

. The effect has been invoked by several authors to explain

elastic anomalies in magnetic materials, usually of the Invar type. Hausch15 criticized the relevance of the Doring effect in most of these cases. Schlosser18 discussed the magnetovolume contribution to the compressibility of Invar, but Hausch15 showed that an exchange - energy contribution is also required in this case. The

present data for iron-chromium -nickel (stainless steel) alloys seem to constitute a clear case of the Doring

effect.

If the Doring effect is responsible for the elastic anomalies observed in stainless -steel-type alloys, then A£^ calculated from Eq. (10) should agree roughly with

the magnitude of the observed anomalies. Apparently,

the volume magnetostriction of these alloys has not yet

been determined. Using data for Invar, 15 an iron alloy

containing about 35v Ni and having the same crystal structure as the alloys considered here, it is calculated from Eq. (10) that A£u /£2 = - 1. 4* 10'" cnr dyn. For

the 304 stainless steel alloy, the present results give for the bulk modulus A£/£2 = - 0. 2 x 10" 14 cmVdyrt.

Thus, the observed anomalies show a reasonable corre- spondence to the magnitude of the predicted Doring effect. A more exact correspondence could be established

if the volume magnetostrictions of these alloys were

known. If the Doring effect is the correct interpretation

of the data, then the volume magnetostriction of these

alloys must be quite large since it is large m Invartype alloys. Besides Invar and other iron-nickel alloys,

large magnetovolume effects have also been observed in iron-manganese, iron-palladium, iron -platinum, and

iron -cobalt alloys. 19 A quantitative correspondence is also precluded because the alloys reported on here were

studied in mechanically deformed states. Koster1, showed that the A£ effect depends sensitively on metallurgical variables. The usual A£ = A£X effect is generally suppressed by mechanical deformation. No studies seem to have been made on the effect of mechanical deformation on A£^. It would be expected that the A£u effect will be enhanced in annealed materials since the

3859 J. Appl. Phys., Vol. 46, No. 9, September 1975

residual stresses due to mechanical deformation would

not interfere with domain -wall motions.

Of the four alloys, the smallest elastic anomalies were observed in the 310 alloy and in the A286 alloy.

These materials have a much higher nickel content. Thus, it is suggested that in stainless -steel -type alloys

the effect of nickel is to reduce the volume magnetostriction but not to suppress the occurrence of the anti- ferromagnetic phase.

The desirability of further studies on these alloys,

especially with respect to the Doring effect, is indi- cated. Ideally, carefully prepared alloy single crystals

would be tested in magnetic fields.

Since changes in elastic constants are'usually accompanied by significant changes in internal friction, it would also be interesting to study the low -temperature

anelastic properties of these alloys. The magnitude of the mechanical damping determines the magnitude of stresses established in vibrating parts. And all of these

alloys are candidate materials for low -temperature uses where vibrations may occur. Magnetomechanical hysteresis has already been established as an important damping mechanism in ferromagnetic materials. 20

V. CONCLUSIONS

From the results of the present study the following

conclusions are drawn:

(i) Stainless -steel alloys AISI 304, AISI 310, AISI 316,

and A286 have qualitatively similar elastic -property

variations with temperature.

(ii) The shear modulus behaves regularly over the entire temperature range studied.

(iii) The elastic constants with dilatational components —Young's modulus, the longitudinal modulus, the bulk modulus, and Poisson's ratio —behave regularly

above about 75 K but anomalously at lower temperatures.

The anomalies are largest for the bulk modulus and

smallest for Young's modulus.

(iv) The anomalies are largest in the AISI 304 and

AISI 316 alloys and smallest in the AISI 310 and A286 alloys. A relationship between the magnitude of the anomaly and nickel content is suggested.

(v) As suggested first by Doring, the anomalies can

be interpreted thermodynamically as the difference be- tween constant -magnetization and constant -field bulk

moduli.

Noted added in proof. Subsequent studies on other samples of these materials showed that the values of £, G, and B reported here for AISI 316 are all higher than the average values by about 6%. All data reported for other

materials are believed to be typical.

ACKNOWLEDGMENT

This work was supported in part by the Advanced

Research Projects Agency of the Department of Defense.

Ledbetter, Weston, and Naimon 3859

109
+NRC-NBS Postdoctoral Research Associate, 1973-4.

t Present address: Rockwell International, Rocky Flats Plant, Golden, Colo. 80401.

'J.J. Gilman, Aust. J. Phys. 13, 327 0 960).

2 This trade name is used to describe the material; its use

does not imply an endorsement by NBS of a particular

product.

! H.J. McSkimin, J. Acoust. Soc. Am. 33, 12 (1961).

4 E.R. Naimon, W.F. Weston, andH.M. Ledbetter, Cryogenics 14, 246 (1974).

5 Y.P. Varshni, Phys. Rev. B 2, 3952 (1970).

6R.P. Mikesell and R. P. Reed, J. Res. Natl. Bur. Std. (U.S.) 70C, 207 (1966). 'P.E.Armstrong and D.T. Eash, in Advances in Cryogenic

Engineering, Vol. 14, edited by K.D. Timmerhaus (Plenum,

New York, 1959), p. 64.

8 E.I. Kondorsky and V. L. Sedov, Sov. Phys. -JETP 35, 1104

(1959).

9W.F. Schlosser, J. Phys. Chem. Solids 32, 939 (1971).

10J.W. Christian, The Theory of Transformations in Metals in Metals and Alloys (Pergamon, Oxford, 1965), p. 214. UW.F. Brown, Jr., Phys. Rev. 50, 1165 0936).

12R.M. Bozorth, W. P. Mason, and H.J. McSkimin, Bell Syst. Tech. J. 30, 970 (1951).

13 0. Engler, Ann. Phys. (Leipz.) 31, 145 0938).

14 W. Doring, Ann. Phys. (Leipz.) 32, 465 (1938).

15 G.Hausch, Phys. Status Solidi A 15, 501 (1973).

,6 E.P. Wohlfarth, Phys. Status Solidi A 10, K39 (1972).

17W. K6'ster, Z. Metallkde. 35, 194 0943).

18W.F. Schlosser, Phys. Status Solidi A 18, 235 (1973).

19W.F. Schlosser, Int. J. Magn. 2, 167 (1972).

20A. Cochardt, Trans. Am. Soc. Mech. Eng. 75, A196 (1953).

3860 J. Appl. Phys., Vol. 46, No. 9, September 1975 110

Ledbetter, Weston, and Naimon 3860
NBSIR

SEMI-ANNUAL REPORT ON MATERIALS RESEARCH

IN SUPPORT OF SUPERCONDUCTING MACHINERY

FATIGUE RESISTANCE AND FRACTURE TOUGHNESS OF

ENGINEERING MATERIALS AT CRYOGENIC TEMPERATURES

R. L. Tobler, D. T. Read, and R. P. Reed

Cryogenics Division

Institute for Basic Standards

National Bureau of Standards

Boulder, Colorado 80302

October 1975

111
Summary: Fatigue and Fracture Toughness

This section of the report contains two manuscripts, the results of which are summarized as follows:

1) The fatigue life of a uniaxial glass-filament reinforced epoxy composite at liquid

helium temperature can be an order of magnitude greater than at room temperature.

Tensile and fatigue results for this material indicate that such composites can be

attractive structural materials for cryogenic applications.

2) The yield strength of an Fe-21Cr-6Ni-9Mn stainless steel tripled between 295 and

4 K, reaching a value of 1.24GNm-

^ (180ksi) at 4 K. Over this temperature interval

the fracture toughness was considerable, and inversely related to yield strength. This

alloy should be useful at temperatures as low as 4 K, for applications requiring a balance of high strength and fracture toughness.

112
FATIGUE RESISTANCE OF A UNIAXIAL S-GLASS/EPOXY

COMPOSITE AT ROOM AND LIQUID

HELIUM TEMPERATURES

Ralph L. Tobler and David T. Read*

Cryogenics Division

Institute for Basic Standards

National Bureau of Standards

Boulder, Colorado 80302

Abstract

Tension-tension axial fatigue tests of a uniaxial glass filamentreinforced epoxy were conducted at 295 K and 4 K. The fatigue life

was found to be an order of magnitude greater at 4 K than at 295 K.

These results are believed to be the first 4 K fatigue data reported

for a composite material.

Key words: Cryogenics; fatigue; fiber composites; liquid helium;

low temperature tests; mechanical properties.

NRC-NBS Postdoctoral Research Associate, 1974-75.

113
FATIGUE RESISTANCE OF A UNIAXIAL S-GLASS/EPOXY

COMPOSITE AT ROOM AND LIQUID

HELIUM TEMPERATURES

Ralph L. Tobler and David T. Read

INTRODUCTION

Glass-reinforced epoxy composites are prospective structural materials

at cryogenic temperatures where high strength-to-density and strengthto-thermal-conductivity ratios are required and where relatively low

elastic moduli can be tolerated. At temperatures from 295 to 4 K, these

composites offer a combination of physical and mechanical properties

» unmatched by structural metals. Unfortunately, structural design data

for Composites at low temperatures are scarce, a fact which hinders

the utility of these materials.

The advent of superconducting machinery prompted the need for

mechanical property data for materials, in a liquid helium environment at

4 K. Accordingly, programs were instituted to establish a data base

[1], Kasen's review [1,2] of the mechanical behavior of glass-reinforced

composites concludes that tensile and fatigue strengths generally increase

between 295 and 76 K; but, on further cooling to 20 K, there is a high

probability that strengths will decrease. According to Kasen [1,2],

below 20 K, tensile behavior has rarely been investigated, and fatigue

data are non-existent.

In this study, fatigue tests of a filamentary-reinforced composite at

room temperature and 4 K are described. Cycle life as a function of applied

stress is evaluated to provide baseline fatigue data for a specific primary

load carrying component of a superconducting motor. Some experimental aspects

of fatigue testing at 4 K are discussed.

114
EXPERIMENTAL PROCEDURE

The material tested was a unidirectional composite of Owens -Corning

S-901* glass roving and an epoxy resin, SCI REZ 80*. Specimens were fabricated

in the form of 12.38 cm long bands, as shown in Figure 1, using a single

material lot and constant fabrication procedure. The specimen is a scaled

down version of 68.5 cm long support bands designed for suspension of superconducting motor components in liquid helium. All specimen dimensions, except

length, are equal to those of the service component. As specified by the

manufacturer, the specimen fiber content is 82% by weight, 69% by volume; the

_3 void content is less than 1% by volume, and the density is 2.1 g

• cm at

room temperature.

Using flanged bushings, the specimens were pinned to the grips of a

cryostat enabling fatigue tests in a liquid helium environment, as shown

in Figure 2. The cryostat frame has a load carrying capability of 100 kN;

it consists of two tubular compression members which are bridged at the

bottom. The lower halves of these members are AISI 304 stainless steel

having a wall thickness of 3.18 mm. The upper sections, which experience

a thermal gradient during testing, were fabricated from fiberglass-reinforced

plastic tubes, 6.35 mm thick. This lightweight construction provides high

specific strength and low thermal conductivity, with some sacrifice of

rigidity; the stiffness of the frame and load train at 4 K is approximately

23 kN/mm. Details of this cryostat design were described by Fowlkes and

Tobler [3].

* Tradenames are used in this report for clarity; in no way does this

imply recommendation or endorsement by the National Bureau of Standards,

nor does it imply that the material is the best available for the

purpose.

115
Room temperature tests were performed with the apparatus in ambient air

at 295 ± 2 K, and at a relative humidity of 30 ± 5%. The tests at 4 K were

accomplished using the double-dewar arrangement illustrated in Figure 2.

The specimen was enclosed in a fiberglass-reinforced plastic dewar, having

a volumetric capacity of 9 liters. This dewar is vacuum insulated, and is

surrounded by a f iberglass-epoxy dewar containing liquid nitrogen. Tests

began by cooling the specimen and apparatus to 4 K, according to the two and

one-half hour cool-down procedure previously described [3]. The liquid level

was maintained several inches above the specimen, and was constantly monitored

with a carbon-resistor level indicator. Liquid helium was continuously

transferred into the cryostat to replenish the amounts evaporated during

fatigue tests.

All tests were conducted using a 100 kN closed-loop servo-hydraulic test

machine. Replicate tensile tests were conducted at a loading rate of

5.7 N • s to determine the nominal fracture strength and load-deflection

(P-6) characteristics of virgin specimens. Load cell and LVDT outputs

were used in recording load-versus-actuator displacement curves. The

nominal fracture strength, a^, was calculated from the maximum load,

P , and the mean cross-section area, 2A, of the unloaded specimen at room

max r

temperature:

a

f = P /2A (1) max

The specimen stiffness, P/6 , was obtained by deducting the load frame

deflection, 6 , from the total actuator displacement, 6 :

P P

6 (2)

S

116
where the load frame stiffness was measured separately by replacing the

specimen with a rigid steel block. The nominal or effective value of

specimen modulus was also calculated:

M

s - irs

< 3>

where L is the initial specimen- length at room temperature, and 6^ is the

specimen deflection at a 10 kN load.

The axial fatigue tests were conducted under controlled load, using a

sinusoidal load cycle, at frequencies from 25 to 27 Hz. The ratio, R, of

minimum/maximum load was 0.1, except as noted in the text. A digital indicator

was used to measure the peak loads, which were controlled to within 2% of

the specified values.

RESULTS

A representative load-displacement record obtained at 4 K is shown in

Figure 3a, and static tensile results are listed in Table 1. There was no

measureable change in stiffness or modulus between 295 and 4 K, but tensile

-2

strength increased by 28% above the room temperature value of 133.4 kN • cm

The scatter in tensile strength at 4 K was higher than at room temperature,

with one specimen at 4 K failing at a stress lower than any of the room

temperature values. This anomalous result was not included in the average

value of fracture strength listed in Table 1: the load-displacement record

of the test of this specimen shows a noticeable increment in the displacement

without a corresponding increase in the load at a low value of the load, as

can be seen in Figure 3b.

The room temperature fatigue results are shown in Figure 4, where the

maximum fatigue stress is plotted as a fraction of the static strength. The

117
data trend and scatter are similar to other results, particularly those

reported by Hofer and Olsen [5] for a unidirectional S-99A glass filamentreinforced epoxy. Hofer and Olsen 's data pertain to conventional hour-glass

-2

specimens having a tensile strength of about 162 kN • cm . Although

their material exhibited a higher cycle life capability at a given absolute

stress value, the results for both composites can be normalized on the basis

of differences in static strength as shown in Figure 4. It is also significant

to note that these materials at room temperature do not exhibit a fatigue

limit at cyclic stress levels as low as 20% of the static strength.

The effect of varying the load ratio was investigated briefly at room

temperature where two specimens were fatigued to fracture at R = 0.5. As

shown in Figure 4, the cycle life appeared to improve at the higher load

ratio; further conclusions are not justified in view of the limited data.

The fatigue data obtained at liquid helium and room temperatures are

summarized in Table 2 and Figure 5. These results clearly demonstrate a

superior fatigue resistance at 4 K. Two specimens survived 10^ stress

7 -2

cycles of maximum levels 29.0 and 30.0 x 10 N • m (17.6% of tensile

strength) at 4 K without failure; their tensile strength and effective

modulus after cycling were only slightly less than the average values for

virgin specimens, as shown in Table 2b. These high values of residual

strength and residual modulus indicate that cyclic loading at these levels

at 4 K does not significantly degrade the load-bearing capability of these

4 composite specimens. A third specimen failed after 1.49 x 10 cycles at a

7 -2

stress level of 71.0 x 10 N • m , exceeding the room temperature endurance

by a factor of 30. A final specimen was cycled for specified intervals at

the increasing stress levels listed in Table 2. The results show that the

118
cycle life at 4 K is at least ten times greater than at room temperature

7 -2

for maximum cyclic stresses from 38 to 82 x 10 N • m

Although a reliable S-N curve at 4 K could not be constructed without

more extensive data, it appears that the stress level required for failure

at a specified number of cycles may approach twice the value required at

room temperature. The improvement in fatigue resistance at 4 K exceeds

that which might be predicted based on scaling the fatigue stress levels

to account for the increase in static strength alone.

Fatigue and tensile specimens fractured at 4 K are shown in Figure 6.

Failure invariably occurred at points of tangency where the specimen contacted

the bushings. Fatigue cracks initiated concurrently at all four of the

points of tangency; the cracks penetrated the specimen ligaments along planes

normal to the tensile axis, and delaminations subsequently emanated from

the gripped regions. The delaminations propagated gradually in a direction

parallel to the loading axis, along the entire length of the specimen. When

fatigue damage reached critical proportions the weakest ligament splintered

into filaments, which was the failure mode of tensile specimens, as shown

in Figure 6. The epoxy at the four points of tangency in specimens fractured

in fatigue at both 4 K and 300 K was quite dark, possibly indicating charring.

Temperature had no obvious influence on the appearance of these failures.

Specimen deflection at a constant fatigue load increased progressively,

as shown in Figure 7 for room temperature specimens fatigued at 25.5 x 10^ N

*

The rate of increase of deflection differed by an order of magnitude for these

specimens, as did their fatigue lifetimes, . This suggests that scatter

in cycle life is due to specimen-to-specimen variability in resistance to

119
fatigue damage propagation, as well as initiation. The loss of stiffness

was always most pronounced during the last few cycles to failure, and

extreme adiabatic heating was also noted in the terminal stages of fatigue.

DISCUSSION

Relating laboratory results to service behavior is a practical problem

which was minimized in this study by performing a limited number of tests under

conditions nearly equivalent to those of the service application. The results

must be regarded as lower bounds on the uniaxial tensile properties of the

material itself, since the gripping fixture influenced the results. However,

the gripping fixture is a practical one for service applications, and the

data reported here are directly applicable in design.

It can be concluded from the 4 K results and known trends between 295

and 20 K [2] that the fatigue and tensile strengths of this material are

lower at room temperature than at any cryogenic temperature. In the application

under consideration, the component will experience temperatures ranging from

295 to 4 K. The room temperature properties are therefore limiting, and

should form the basis for design. In certain applications where a component

operates exclusively at cryogenic temperatures, it would also be possible

to design conservatively on the basis of room temperature properties. However,

a maximum level of design efficiency at cryogenic temperatures requires more

extensive low temperature mechanical property data.

In many structural applications, it is essential to monitor fatigue

damage and to replace critical components before catastrophic failure

becomes probable. The occurrence of delaminations in this composite after

120
25% of total cycle life provides obvious evidence of degradation. Damage

could be detected earlier by dismounting the specimen to examine the

surfaces in contact with the bushing. Surface deterioration and discoloration were noted at these locations. The loss of stiffness also provides a

means of detecting fatigue damage. Thus, a simple inspection of service

components should provide ample evidence of degradation prior to failure.

Cyclic stressing produces specimen heating and agitation of the cryogen.

This agitation increases the heat leak into the cryogen due to convection,

because the cryogen is splashed into the warmer regions of the cryostat.

The heat input to the cryogen from both of these sources increases strongly

with displacement. In these tests the helium loss under static load was

only 0.15 I • h . This increased during fatigue tests, varying from 7 to

15 I ' h j over the range of stress levels investigated.

Adiabatic heating in glass-reinforced plastics results from their high

internal friction and low thermal conductivities [6]. Internal temperature

rises as high as 70°F were noted in room temperature fatigue tests of such

materials [7]. At 4 K, adiabatic heating should be more significant since

the thermal conductivity of glass reinforced epoxies is lowered by a factor

of three [4]. Furthermore, the temperature in the interior of the specimen

depends on the specimen thickness and cycling rate. Since the present data

pertain to a specimen thickness and frequency equivalent to service conditions,

the effects of adiabatic heating on fatigue resistance here are academic

from an applications viewpoint. However, this effect must be accounted

for in future tests if fatigue results are to be independent of specimen

geometry and test variables.

121
CONCLUSIONS

From these tests, the following conclusions can be drawn:

1. The tensile strength and fatigue resitance exhibited by a glass

roving reinforced epoxy plastic (GFRP) composite material indicate that

such composites can be attractive structural materials for cryogenic

applications.

2. The fatigue life of a GFRP composite material can be an order

of magnitude greater at liquid helium temperature than at room temperature.

3. A measurable decrease in the stiffness of GFRP specimens may

precede high cycle fatigue fracture at 300 K.

ACKNOWLEDGMENT

The cooperation of M. Calderon of Airesearch, Torrance, California

in providing the specimens and reviewing the manuscript is gratefully

acknowledged.

122
REFERENCES

1. R. P. Reed and A. F. Clark, Eds., "Materials Research for Superconducting

Machinery," National Bureau of Standards, Boulder, Colorado, AD

780596 and AD/A 004 586, Nat. Tech. Inf. Serv., Springfield, Virginia

(1974)

.

2. M. B. Kasen, "Mechanical and Thermal Properties of Filamentary-Reinforced

Structural Composites at Cryogenic Temperatures-I : Glass-Reinforced Composites", Cryogenics, Vol. 15, No. 4, 230-240 (1975).

3. C. W. Fowlkes and R. L. Tobler, "Fracture Testing and Results for a

Ti-6A1-4V alloy at Liquid Helium Temperature," to be published.

4. M. Segimoto, "Qualification Test Results for Glass Filament-Wound

Support Bands," Report No. 7581, Structural Composites Industries,

Inc., California (Jan. 1975).

5. K. E. Hofer, Jr., and E. M. Olsen, "An Investigation of Fatigue and

Creep Properties of Glass Reinforced Plastics for Primary Aircraft

Structures," Report M6104, 111. Inst. Tech. Res. Inst. (Apr. 1967).

6. L. J. Broutman, "Fibre-reinforced Plastics," Chap. 13 in Modern Composite

Materials , L. J. Broutman and R. H. Krock, Eds., Addison-Wesley

,

Massachusetts (1967)

.

7. R. B. Heywood, Present and Potential Fatigue and Creep Strengths of

Reinforced Plastics, Tech. Note. No: Chem. 1337, Royal Aircraft

Establishment, Ministry of Supply, London, W. C. 2 (Oct. 1958).

LIST OF TABLES

Table 1. Static tensile properties of S-901 glass/epoxy specimens.

Table 2a. Fatigue results for S-901 glass/epoxy at 295 K.

Table 2b. Fatigue results for S-901 glass/epoxy at 4 K.

123
Table 1. Static tensile properties of S-901 glass/epoxy specimens.

Mean X-sect. Stiffness Effective Nominal Failure

Temperature Specimen Area, 2

A

Modulus

,

M

s Strength, 0^

(K) (No.) ((cm\) (10

5

N'nT

1

) (10

9

N«m" (10

7

N«m"

2

)

295 l

a

.3612 N/A N/A

if

134.6

2 .3604 ii 1

1

133.8

3a .3648 ii !l 131.6

Avg. JL JJ . H

13 .3677 159 62.8 N/A

14 .3652 148 61.8 it

19 .3613 151 60.5 it Avg. 153 61.7

4 4 .3664 160 66.7 158.8

5 .3664 147 61.1 182.1

21 .3652 145 60.6 126.

7

b

Avg. 151 62.8 170.4

Tests performed by the manufacturer [4]

.

lot included in the average (see text)

.

124
Table 2a. Fatigue results for S-901 glass/epoxy at 295 K.

Maximum Cyclic

Specimen Stress, Fatigue Cycles

7 „ to failure

(No.) (10 N-m ) (No.)

1 o 13 25 .

5

1 . 24b X in6 10

±y 25 .

5

5 . /b X in5

1U

y on q 2y . j b . /5 X in4

1U

lb 2y . j o . 3b X m4

lo on o 29. 3 / . Jo X in4

10

7 31 .

0

1 . 9b X in5 10

12 31.0 1.42 X

5

10

20 31.0 7.46 X io

5

6 37.9 8.85 X io

3

10 37.9 2.22 X 10*

11 37.9 2.09 X 10*

24 71.0 3.93 X io

2

22

a 41.4 2.00 X io4

23

3

55.2 5.40 X io

3

aR = 0.50

125
Table 2b. Fatigue results for S-901 glass/epoxy at 4 K .

Specimen

(No.)

1

Maximum Cyclic

Stress, 0^ Fatigue Cycles Residual Strength

(10

7 N'nf2

) (No.), (10

7 N*m~

2

)

Res'idual Modulus

. (10 N«m

QO zy .u i X xU v.^UI^ OUC J Xjy 58.2

T "7 17 on ft i 31 ,U 1 „ -t rfi II ici: X 1U Ijd D_> . O

14 71.0 1.49 x 10

4 (failure) N/A N/A

15

b - 1 37.9 2.36 x 10"* (run-out) ii

2 44.8 6.92 x 10

4 it

3 55.2 1 x 10

4 ii

4 62.1 1 x 10

4 ii

5 69.0 1

4 x 10 ii

6 82.7 2.52

4 x 10 (failure) ti

b

Cycled at increasing stress levels.

LIST OF FICURES

Figure 1. S-901 glass/epoxy test specimen.

Figure 2. Liquid helium fatigue cryostat.

Figure 3. Static test records; (a) ordinary failure,

(b) premature failure.

Figure 4. Room temperature fatigue data.

Figure 5. 4 K fatigue results.

Figure 6. Tensile and fatigue failures of specimens tested at Ii K.

Figure 7. Def lection-versus-cycles for specimens tested at 295 K.

126
(/> O —

I

<D 1—

i

*3CO

\

\

\

A

I I 1111] I I I I I I ll| 1—I TTT

• - This Study O - Hofer, Olsen [3]

S-901 Glass/Epoxy S-994 Glass/Epoxy

a

f =133x10 Nnf2 af=162xl07 Nm"

?

f = 25-27 Hz f = 0.1-30 Hz

R = 0.5

•R = 0.1

Mean, R = 0.05

c

III L 11 I 1 I I M 1 1 1 I 10- 10" 10 3

FATIGUE CYCLES, N

FIGURE 4 - Room temperature fatigue data.

10 D 10

TTTTT TT

\

\

\

TT

S-901 Glass/Epoxy • 4K O Specimen "15, 4K — No Failure

R = 0.1, f = 25-27 Hz

295K, Mean

Uil

10

J 10" io

J FATIGUE CYCLES, N

10

u 10

FIGURE 5 - 4K Fatigue results.
Figure 6. Tensile (lower) and fatigue (upper) failure of specimens

tested at 4 K.

FATIGUE CYCLES. N

FIGURE 7 - Deflection-versus-cycles for specimens at 295K.

130
TENSILE AND FRACTURE BEHAVIOR OF A NITROGEN-STRENGTHENED,

CHRONIUM-NICKEL-MANGANESE STAINLESS STEEL AT CRYOGENIC TEMPERATURES*

R. L. Tobler and R. P. Reed

Cryogenics Division

Institute for Basic Standards

National Bureau of Standards

Boulder, Colorado 80302

ABSTRACT

J-integral fracture and conventional tensile properties are reported

for an Fe-21Cr-6Ni-9Mn austenitic stainless steel which contains 0.28 N as an

interstitial strengthening element. Results at room (295 K) , liquid nitrogen

(76 K) , and liquid helium (4 K) temperatures demonstrated that the yield strength

and fracture toughness of this alloy were inversely related, exhibiting opposite

temperature dependences. Over the temperature range investigated, the yield

8 -2

strength tripled, reaching a value of 12.4 x 10 N*m (180 ksi) at 4 K. The

fracture toughness, as measured using 3.8 cm thick compact specimens, decreased

considerably from 295 to 4 K. During plastic deformation at 295 K the alloy

undergoes little martensitic phase transformation, but at 76 and 4 K it transforms extensively to hep and bec martensitic products. The amount of bec

transformation product was measured as a function of elongation.

Key words: Cryogenics; fracture; low temperature tests; martensitic

transformations; mechanical properties; stainless steel alloys.

NBS contribution, not subject to copyright.

131
Introduction

Recently, austenitic stainless steel strengths have been increased considerably

by the substitution of nitrogen and manganese for nickel. In addition to the

interstitial and solid solution strengthening, these elements serve to retain, and

in some cases, increase the austenite stability with respect to martensitic transformations. Compared to nickel, these elements are more abundant and less expensive.

The alloy studied in this report, Fe-21Cr-6Ni-9Mn-0. 3N (21-6-9) , has a room

temperature yield strength nearly twice that of AISI 304. Available tensile and

(1-4) -j i v impact data suggest that the alloy retains good toughness at low temperatures,

leading to consideration of its use for applications benefiting from high strength

and toughness.

Accordingly, 21-6-9 is currently being considered for such critical components

as the coil for the prototype controlled thermonuclear reaction superconducting

magnets and the torque tube for rotating superconducting machinery. To insure

satisfactory service life and to compare with other material contenders it is

necessary to evaluate the fracture resistance of the alloy. This study presents

the first fracture toughness data for this alloy.

Material

The 21-6-9 austenitic stainless steel was provided by Lawrence Livermore

Laboratories, Livermore, California. The chemical composition (in weight percent)

of this heat is 19.75 Cr, 7.16 Ni, 9.46 Mn, 0.019 C, 0.15 Si, 0.004 P, 0.003 S,

and 0.28 N.

132
The steel was subsequently hot rolled from 30.5 x 30.5 x 10 cm slabs to

50 x 50 x 3.6 cm plate using the following schedule:

soak for 4 hours at 1366 K;

roll from 10 cm to 9.4 i zm, screw down (compress) to 9.3 cm;

roll from 9.3 cm to 8.7 cm, screw down to 8. 6 cm, rotate 90°;

roll from 8.6 cm to 8.0 cm, screw down to 7. 9 cm;

roll from 7.9 cm to 7.4 cm, screw down to 7. 3 cm, rotate 90°;

roll from 7.3 cm to 6.9 cm, screw down to 6. 7 cm;

roll from 6.7 cm to 6.3 cm, screw down to 6. 2 cm, rotate 90°;

roll from 6.2 cm to 5.9 cm, screw down to 5. 7 cm;

roll from 5.7 cm to 5.4 cm, screw down to 5. 3 cm, rotate 90°;

roll from 5.3 cm to 5.G cm, screw down to 4. 8 cm;

roll from 4.8 cm to 4.5 cm, screw down to 4. 3 cm, rotate 90°;

roll from 4.3 cm to 4.2 cm, screw down to 4. 0 cm;

roll from 4.0 cm to 3.8 cm, screw down to 3. 6 cm.

temperature after this hot rolling was 1089 K. Each plate was

then annealed at 1283 K for 1-1/2 hours and air cooled, followed by an anneal

at 1366 K for 1-1/2 hours and a water quench. The resultant hardness was

Rockwell B92 and the average grain diameter was 0.16 mm.

Procedure

Tensile

:

Tensile specimens were machined following ASTM specification E8-72 ; the

reduced section diameter was 0.635 cm and length was 3.8 cm.

The tensile axis was oriented transverse to the final rolling direction of the

plate. Tests were performed using a 44.5 kN screw-driven machine at a cross-

( 6

)

head rate of 0.05 cm per minute. The cryostat assembly was designed by Reed.

133
Measurements at 76 K used liquid nitrogen and at 4 K used liquid helium. Load

was monitored with a 44.5 KN commercial load cell while specimen strain was

measured with a clip-on, double beam, strain gage extensometer . Yield strength

was determined as the flow stress at 0.2% offset plastic strain.

Magnetic

:

To detect the amount of ferromagnetic, body-centered cubic (bcc) martensitic

phase in the paramagnetic, fee austenitic matrix a simple bar-magnet, torsion

balance was used. Previous measurements on Fe-Cr-Ni austenitic steels established a correlation between the force required to detach the magnet from the

specimen and the percent bcc martensite. This same correlation was used for

this study to estimate the amount of bcc martensite in the Fe-Cr-Ni-Mn alloy.

Fracture:

The J-integral specimens were 3.78 cm thick compact specimens of a geometry

(8) described in ASTM E-399-74. The specimen width, W, and width-to-thickness

ratio, W/B, were 7.6 cm and 2.0, respectively. Other dimensions are shown in

Figure 1. The notch, machined parallel to the final rolling direction of the

plate, was modified to enable clip gage attachment in the loadline.

The J-integral test specimens were precracked using a 100 kN fatigue test-

(9) ing machine and attached cryostat. Specimens were precracked at their test

temperature. All precracking operations were conducted using load control and

a sinusoidal load cycle at the rate of 20 Hz. Maximum fatigue precracking loads

(P ) were, except for room temperature, well below the maximum load to fracture

after precracking (P ) as indicated in Table 1. The maximum stress intensity max

during precracking (K^) and the approximate final average relative crack length

(the ratio of the crack length, a, to the specimen width, W) are also listed in

Table 1.

134
Subsequent to precracking, the specimens were transferred to a 267 kN

(60,000 lb) hydraulic tensile machine for fracture testing. Thus, the 76 and 4 K

fracture specimens were warmed to room temperature prior to testing at 76 and 4 K.

This was necessary since the load limitations of the fatigue machine precluded

loading this alloy to fracture at the low temperatures.

The J-integral tests followed a resistance curve technique similar to that

described originally by Begley and Landes . ^

' ^ A series of nearly identical

specimens was tested at each temperature. Each specimen was loaded to produce

given amounts of crack extension. The specimens were then unloaded and heat

tinted to mark the amount of crack extension associated with a particular value

of J. The amount of crack extension could be identified after fracturing the

specimen into halves.

(12) Using the approximate solution for deeply cracked compact specimens,

the value of J for each test was calculated from the total area, A, under the

load-versus-loadline deflection record. The five values of J obtained at each

temperature were plotted versus the average crack extension, la, which was

measured at five locations equidistant across the specimen thickness, and

averaged

.

The critical value of the J integral, J^, identified as the J value at the

initiation of crack extension, is obtained by extrapolation of the best fit J-Aa

curve to the point of actual material separation. Conversion to the plane-

(10-12)

strain fracture toughness parameter, K , is made using

J = 2A/B(W-a)

,

(1)

K

2 E

1-v

(2)

TC

where E is Young's modulus and V is Poisson ratio. Here E is taken as

135
the range of the ratio of the length of the crack near the specimen edges

(a ) to the length of the cracks near the specimen centers (a ) is included in

e c Table I. To satisfy the fracture criteria of ASTM E-399-74 ^ a value of

a

e

/ a

c = 0.90 or greater is necessary.

Table 1. Precracking Parameters for Fracture Test Specimens

Test P^/P x 100 K_ Final Average

f max f &

Temperature ,

. ,^ n 6 „ 2 1/2 S (percent) (10 N»m *m ) Relative Crack a /a

(K) Length, a/W 6 C

295 40-45 48-54 0.638 0.90-0.91

76 22-27 52-63 0.640 0.87-0.90

4 30-35 52-63 0.647 0.88-0.89

Results and Discussion

Tensile:

The yield and tensile strengths, elongation, and reduction of area were obtained for the 21-6-9 alloy at 295, 76, and 4 K. These data are summarized in

Table 2. The data from this study are combined in Figures 2-4 with the unpublished results of Landon^^ for the same heat, also hot rolled and annealed, and

with the results of Scardigno,

^

Malin, , and Masteller^ on annealed bar

stock. The correlations are very good except for the ultimate strength data of

Masteller. Although his alloy was described as annealed, perhaps it was tested

in a worked condition. The spread of the Malin data represent specimens machined

from both longitudinal and transverse directions to the rolling direction.

Typical stress-strain curves at each temperature are presented in Figure 6. The

large discontinuous yield behavior at 4 K probably is associated with adiabatic

specimen heating. Significant local heating is indicated, as the flow stress

136
drops to stress levels less than that which is sustained at 76 K. An

indication of significant local heating is the rise of the reduction of area

to values higher than that obtained during 76 K tests. Specimens tested at

4 K developed very local areas of increased plastic deformation, which resulted

in sizable specimen necking prior to fracture. These load drops should not

be attributed to martensitic phase transformations for three reasons: (1) more

extensive transformation was detected in this alloy at 76 K, compared to 4 K

(see later discussion) and no discontinuities in the stress-strain mode at 76 K

were observed, (2) load drops have been observed in both metastable (e.g.,

AISI/304) and stable (e.g., AISI 310) austenitic stainless steels at 4 K and

(13) no distinction is apparent between the two alloy groups, and (3) in austenitic

steels the amplitude and frequency of the load drops at 4K is a function of the

(13) strain rate which would be expected if local heating were responsible.

From these data it is clear that the alloy 21-6-9 has a significant decrease

of ductility in the temperature range below 195 K, and that tensile elongation

continues to decrease to 4 K.

A primary advantage offered by this alloy is its higher yield strength with

respect to other austenitic alloys. At room temperature the yield strength of

Q _9

the 21-6-9 alloy is about 3.8 x 10 N«m (55 ksi) , compared to AISI 300 series

8 —2

steel values of 2.1 to 2.5 x 10 N«m (30-35 ksi). The yield strength of the

8 —2

21-6-9 steel approximately triples to a value of 12.4 x 10 N*m (180 ksi) as

the temperature is decreased to 4 K. The Fe-Cr-Ni austenitic alloys achieve

values about double or triple their room temperature value (60-110 ksi) at 4 K.

The strength advantage offered by 21-6-9 is greatest at lower temperatures.

137
Fracture

The load-versus-loadline deflection curves at 295, 76 and 4 K are

shown in Figure 6. The curves at 295 K extended to larger deflections than

indicated on the axis of the diagram. The crack remained stable and did

not propagate quickly in the J-Aa tests of room temperature tests. The fracture

test data are tabulated in Table 3. There are no ASTM E-399-74 valid K data;

the 5% secant offset data are denoted K„. Both the ASTM E-399-74 thickness and

Q

crack front curvature criteria are not satisfied at the temperatures 4, 76 and

2 295 K. Using B > 2.5 (K^/a ) , a specimen thickness of 4.2 cm at 4 K is required,

slightly larger than the 3.8 cm thickness. As shown in Figure 7, the crack

front curvature is slightly excessive as the crack lengths of the surface are between

88 and 89% of the average of the internal crack length; 90% is suggested in E399 as the minimum deviation.

The J versus Aa results at room temperature are plotted in Figure 8. Large

extensions were observed due to plastic deformation at the crack tip; only in

two specimens at the highest values of Aa was actual material separation noted.

These two values fall on the same linear curve as the specimen data that did not

exhibit material separation and, furthermore, the entire J versus Aa linear

curve has approximately the same slope as that suggested (J/2a^) by Landes and

Begley?"^' It seems that the response of this extremely ductile material to J

integral tests at room temperature is inconclusive, with no well-defined break

from the linear portion of the plastic deformation curve observable.

According to the tentative size criterion suggested by Landes and Begley^"^ ,

the specimen thickness for valid measurements should satisfy the

relationship

:

B > a(J/a

f

)

where a is a factor greater than 25 and represents the average of the,

yield and tensile strengths. In the tests at 295 K, the conditional

138
2 2 critical J values were in the range 1350 to 1450 x 10 N • m/m . Using

8 —2

the flow stress value of 5.3 x 10 N • m (76.5 ksi) , the J-integral results

at room temperature are invalid for the specimen thickness tested here.

Apparently, a specimen thickness of at least 6.3 cm (2.5 inches) is needed

to insure valid data.

The J-resistance curve at 76 K is also shown in Figure 8. The data fit

a regular trend, with the exception of the datum point representing the

largest observed crack extension. The curve drawn for the remaining

specimens indicates that crack extension initiates at J values in the

range 310 < J < 350 x 10

2.2N'm/m . The corresponding value of K^

C (J),

1/2 2 1/2 estimated using Equation 2, is 318 N • m /m + 5% (292 ksi in ).

At 4 K, the alloy approached linear-elastic behavior, but the results

(8)

of the first three tests failed to satisfy the ASTM validity criteria for

direct measurements. These data are listed in Table 3. Three J-Aa tests

were conducted and these are included in Figure 8. Although more thorough testing

is needed to confirm these, it appears that the J-Aa curve is horizontal. The

3 2

resulting J is about 150 x 10 X*m/m .

Phase Transformations

Subsequent to tensile tests at 76 and 4 K, it was noticed that the

deformed specimens were magnetic. Therefore, these specimens were measured,

using bar-magnet, torsion balance equipment, to correlate magnetic attraction

with specimen reduction of area. The magnetic readings were converted to

percent bcc martensite and the reduction of area converted to elongation, assuming

constant volume. These data are plotted in Figure 9 and typical microstructures

are shown in Figure 10. Several aspects deserve discussion.

Although not positively identified, it is extremely probable that hep

martensite has formed in the 21-6-9 alloy during low temperature deformation.

139
The microphotographs after deformation at 4 K (Figure 10) identify transformed

regions which are parallel to the (111) slip band traces. These appear identical

to the hep areas identified in earlier research on AISI 304, an Fe-Cr-Ni alloy.

The amount of bec martensite formed is large and only slightly less than

that which is formed in AISI 304 at the same temperatures^ Permeability

values of the order of 10 were measured in heavily deformed specimen portions

at 76 K. It is difficult to identify bec martensite in the Figure 10 photomicrographs. Normally, in austenitic stainless steels the bec martensitic

product has an acicular, plate-like morphology with the habit plane of the

plate not {ill}. Examination of specimen microstructures, typified by Figure 10

indicate that only at {ill} band intersections are plate-like and distinctive

microstructures observed.

There is clear evidence that the amount of the transformation is

suppressed, as a function of either stress or strain, at 4 K when compared

to 76 K measurements. This is similar to the Fe-Cr-Ni (AISI 304) alloy

martensitic transformation behavior ^'^"^ . Our earlier work^'^""^ also

indicated that the hep martensitic phase was suppressed at 20 and 4 K.

Apparently, in the complicated energy balance affecting martensitic transformation for these alloy systems at low temperatures the increase of flow

stress and the decrease of dislocation mobility more than offset the gradually

increasing free energy difference between the structures.

Finally, it is not clear that the observance of martensitic transformations is deleterious to material application. Normally, the material is

used in service at stress levels less than the yield strength; under these

conditions no martensitic transformations occur. The complexities and concern

usually are discussed when one considers welds and weld techniques. Chemical

140
segregation and stress concentrations are both more likely, permitting particular

sections to be less stable and, locally, stressed above the yield strength. In

these situations martensitic products will form.

AISI 304 behaves in a similar manner; it is stable on cooling to low

temperatures but transforms to hep and bec martensitic products during plastic

deformation. But, unlike 21-6-9, the fracture toughness of AISI 304 remains

(14) extremely high at 4 K This implies that martensitic transformations are not

harmful to the fracture toughness of the Fe-Cr-Ni stainless steel. This is not

clear in the case of the Fe-Cr-Ni-Mn-N alloy, however, as the toughness is

rapidly decreasing at 4 K. For appropriate safety of operation at 4 K, additional

research is necessary to understand the effect of martensitic transformations on

fracture toughness.

Acknowledgement s

The authors would like to thank Dr. P. R. Landon, Lawrence Livermore

Laboratories, for kindly supplying the rolled material. Dr. R. P. Mikesell

conducted the tensile tests, R. L. Durcholz contributed technical assistance to

tensile, fracture, and metallographic preparation, and Dr. M. B. Kasen was

responsible for the photomicrographs.

141
References

1. P. R. Landon, Unpublished data, Lawrence Livermore Laboratories,

Livermore, CA (1975)

.

2. P. F. Scardigno, Masters Thesis, Naval Postgraduate School, Monterey,

CA, AD/A-004555 (1974). r

3. C. 0. Malin, NASA SP-5921 (01) , Technology Utilization Office, NASA,

Washington, D.C. (1970).

4. R. D. Masteller, NASA CR-72638 (N70-27114) , Martin Marietta Corp.,

Denver, CO (1970).

5. "Tension Testing of Metallic Materials", E8-72, Annual Book of ASTM

Standards, Part 31 (ASTM, Philadelphia, PA, 1973).

6. R. P. Reed, Advances in Cryogenic Engineering , Vol. 7 (Plenum Press,

New York, 1962) p. 448.

7. R. P. Reed and C. J. Guntner, Trans. AIME, 230 , 1713 (1964).

8= ASTM Standards, E399-74, part 10 (ASTM, Philadelphia, PA 1974).

9. R. P. Reed, R. L. Tobler, and R. ,P. Mikesell, "Advances in Cryogenic

Engineering - ICMC" , Vol. _22 (Plenum Press, New York, 1976), to be

published. s

10. J. D. Landes and J. A. Begley, Westinghouse Research Laboratory

Scientific Paper 73-1E7-FMPWR-P3, Pittsburgh, PA (1973).

11. J. D. Landes and J. A. Begley, "Fracture Analyses", ASTM S TP 560,

(American Society for Testing and Materials, Philadelphia, PA 1974)

p. 170-186.

12. J. R.Rice, P. C. Paris, and J. G. Merkle, ASTM STP 536 (ASTM,

Philadelphia, PA 1973) p. 231.

13. C. J. Guntner and R. P. Reed^ ASM Trans. 55, 399 (1962).

14. R. P. Reed, A. F. Clark and E. \C. van Reuth, Eds. "Materials Research for

Superconducting Machinery III", AD-A012365/3WM, National Technical

Information Service, Springfield, VA (1975).

15. R. R. Vandervoort, Metals Engr. Quart. 12, No. 1, 10 (1972).

\

\

142
List of Figures

Figure 1. Compact specimen for fracture testing of Fe -19Cr-6Ni-9Mn

alloy.

Figure 2. Summary of tensile and yield strength data as a function of - temperature for the Fe-21Cr-6Ni-9Mn alloy.

Figure 3. Summary of tensile elongation as a function of temperature for

the Fe-21Cr-6Ni-9Mn alloy.

Figure 4. Summary of tensile reduction of area as a function of temperature

for the Fe-21Cr-6Ni-9Mn alloy.

Figure 5. Stress-strain curves for the Fe-21Cr-6Ni-9Mn alloy at

295, 76 and 4 K.

Figure 6. Typical load-deflection curves at 295, 76 and 4 K for annealed

Fe-21Cr-6Ni-9Mn alloy.

Figure 7. Fracture characteristics of Fe-21Cr-6Ni-9Mn alloy at

295, 76 and 4 K.

Figure 8. The J-integral as a function of crack extension at 295, 76 and 4 K

for annealed alloy Fe-21Cr-6Ni-9Mn.

Figure 9. Estimated percent bcc martensite that forms during tensile tests

as a function of tensile elongation.

Figure 10. Microstructures of alloy 21-6-9 after deformation at 4 K. Bands

lie on {ill; austenite planes and probably represent hep and bcc

martensite, (a) 800X, (b) 1200X.

List of Tables

Table 1. Precracking Parameters for Fracture Test Specimens.

Table 2. Tensile Properties of Fe-21Cr-6Ni-9Mn alloy.

Table 3. Fracture Results for 377 cm Thick Compact Tensile Specimens

of Alloy Fe-21Cr-6Ni-9Mn.

143
Table 2. Tensile Properties of Fe-21Cr-6Ni-9Mn Alloy

Yield Strength Tensile Strength Elongation

0.2% offset

(10

8

N-m"

2

)

Temperature 0.2% offset n n

~^"\ 2.5 cm gage Reduction of Area

(K) /in8 -2,

U ' length (percent)

295 K 3.50

3.57

Average 3.53 (51 ksi)

76 K 9.13

8.86

Average 8.99 (130 ksi)

4 K 12.58

12.24

(percent)

6.96 61 79

7.05 61 78

7.01 (102 ksi) 61 78

14.62 42 32

14.85 43 41

14.74 (214 ksi) A3 37

16.33 16 40

16.34

Average 12.41 (180 ksi) 16.34 (237 ksi) 16 40

144
Table 3.

Temperature a/W J Aa

Q 3/2

\ MN/ m J (K)

3 ?

V-lu in m/ m ^ (cm)

295 0.638 177 0.013* 58

0U»OjU fi^fi 7/i / U. Uj1 x bl

0 fiAf) yto _>_> UiDJJ J.JJJ u . uy

/

K O J n ka? 1 A9 r\ 119* jU

76 0.612 261 0 134

0.634 413 0.028 153

U • O'+U nyy 131

A0/47/, u. u/y iI nJ / 788 u . uyi Ton 13U

u . oh j 698 n i no u . iyo 130

/, u . o / u U

0.655 147 0.020 167

0.670

III?]

173

0.080 158

0.645 N/A 164

0.648 N/A 162

U. o43 n/a 159

0.656 162 .076

^ in* lb

in

= 1.75 x 10

2 N **

1

ks±

= 1.099

2 in m

x 10 —

2

m

-1/2

•m

( a ) from E-399-74 method; invalid due to insufficient thickness.

k

calculated at the first load discontinuity (pop-in)

(c) calculated at the maximum load.

* refer not to crack opening distances, but to stretch zone distances

145
9.1cm

0.79cm Diet.

1.9cm Dia.

< l ~r lem 1.9cm

Thickness (B) = 3.8cm

Figure 1. Compact specimen for fracture testing of Fe-19Cr-6'Ni-9Mn alloy.

I I I I I I

I I I I I I J

0 50 100 150 200 250 300

TEMPERATURE, K

figure 2. Summary oC tenaile unci yield .itrcngtli dahu d.3 a function of temperature for tHe Fe-21Cr-feNi-9Mh alloy.
Figure 4. Summary of tensile reduction of area as a function

of temperature for the Fe-21Cr-6Ni-9Mn alloy.

80

60

u 40

20

——

REDUCTION of AREA

O This Report

J • London \ly/ b)

I Malin (1970)

1

I ! 1 ! ! 1 1 1

80

60

u

ui 40

Q_

20

50 100 150 200

TEMPERATURE, K

250 300

ELONGATION

V Masteller (1970)

0 This Report

A Scardigno (1974)

London (1975)

1 Malin (1970)

50 100 150 200

TEMPERATURE, K

250 300

figure 5. Summary of tensile elongation as a fund ion oi temperature for the Fe-2 1 Cr-b.N i-9Mn alloy...

149
Sox QUO U-i D. O

r-H O

t a o

3 C

4J O O

3 QJ ^)

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O

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311SN11HVW "q lN33H3d Q31VWI1S3

150
NBSIR

SEMI-ANNUAL REPORT ON MATERIALS RESEARCH

IN SUPPORT OF SUPERCONDUCTING MACHINERY

MAGNETOTHERMAL CONDUCTIVITY

L. L. Sparks

Cryogenics Division

Institute for Basic Standards

National Bureau of Standards

Boulder, Colorado 80302

October 1975
Summary; Magneto thermal Conductivity

The thermal conductivity and effect of a magnetic field on the thermal conductivity

of S 304 00 have been determined in the temperature range from 5 to 20 K. The results

indicate that a 6366 kA/m (80 kOe) magnetic field reduces the thermal conductivity of this

material by 9.2% at 5.5 K and by 6.7% at 19.75 K. The present data at H = 0 compares very

well with literature data on similar materials. Several minor modifications have been

incorporated primarily to provide better thermal contact between the specimen and the

thermometers. Based on the S 304 00 tests, the precision of the measurements has been

improved. The reproducibility of the results now appears to be about 5% in thermal con- ductivity except at the temperature extremes where it is near 8%.

Contents: Magneto thermal Conductivity

Page

1. Introduction 153

2. Procedures 153

2.1 Apparatus 153

2.2 Materials 154

3. Results and Discussion 154

4. Conclusions 155

5. References 156

List of figures 156

Preprint: Magneto thermal Conductivity of Selected Pure Metals and Alloys . . . 160

152

i

|

i
MAGNETOTHERMAL CONDUCTIVITY

L. L. Sparks

Cryogenics Division

Institute for Basic Standards

National Bureau of Standards

Boulder, Colorado 80302

1 . Introduction

The objective is to determine the effect of magnetic fields on the thermal conductivity of technically important metals. The need for this information arises from the

development of rotating machinery operating at cryogenic temperatures. The existing world's

literature on magnetothermal conductivity, X(H), is concerned almost exclusively with

scientific materials, e.g., very pure materials and single crystals. A complete bibliography of the subject was given by Sparks and Fickett [1].

The materials studied in this program are being used or are candidates for use in superconducting motors and generators. Optimum design of these machines, which must operate at lov, temperatures while in magnetic fields, requires a detailed knowledge of how the thermal properties of the constituent materials are affected by a magnetic field. The broad material

categories of interest include superconductor stabilizing materials such as copper and aluminum, and structural materials such as nickel alloys, stainless steels, and metallic composites

Results of the \(H) work were presented at three conferences since the last semi-annual

ARPA report: American Physical Society (Denver, March), 14th International Conference on Thermal Conductivity (University of Connecticut, June), and the joint meeting of the Cryogenic Engineering Conference and the International Cryogenics Materials Conference (Kingston,

Canada, July). A preprint of the paper presented at the ICMC Conference is included at the

end of this report. The work was well received in each case and led to several worthwhile

discussions on the effect of a magnetic field on metals.

Several changes have been made in the apparatus and instrumentation in an attempt to

improve the precision of the results. The details of these changes are discussed in the

procedures section of this report. Tests on S 3040 0 (AISI 304 stainless steel) have been

used to evaluate the system with the changes incorporated.

2. Procedures

2.1 Apparatus

The principle components of the X(H) system are shown schematically in figure 1. The

basic operation of the system has been described in previous reports [1,2] and will not be

repeated here except in relation to the changes.

Thermal conductivity is calculated using the equation

A/£ AT

'

Errors made in the measurement of the temperature difference along the specimen, AT, are reflected directly in X. Perhaps the most difficult problem encountered in measuring small

temperature differences is that of making good thermal contact with the temperatures of

interest. Before the current changes were instituted, the carbon resistance thermometers,

(CRT's), were press fit into wells in the THERMOMETER BLOCK and packed with a thermally

conducting grease. The THERMOMETER BLOCKS were machined so that knife edge contacts were

made when they were clamped to the specimen. This system had the following shortcomings

which are eliminated or minimized by the changes made in the method of sensing the specimen

temperatures: (1) It was impossible to be sure that the thermal contacts between the specimen and the THERMOMETER BLOCK and between the THERMOMETER BLOCK and the CRT's were the same
for the upper and lower THERMOMETER BLOCKS; (2) the THERMOMETER BLOCKS were comparable to

the specimen in mass and heat capacity. This was the case even though the dimensions of

the BLOCKS were as small as possible consistent with the CRT wells and the clamping screw;

(3) it was difficult to attach the clamps to the small specimens without introducing some strain in the specimen. The new "thermometer holder", designed to alleviate the problems

discussed above, is shown schematically in figure 2. Thermal contact to the specimen is

achieved by soldering the 20 AWG copper wire directly to the specimen; this method assures

good thermal contact and should be identical for both the upper and lower thermometer

stations. Thermal contact and electrical isolation between the copper wire and' the CRT is

provided by the sapphire chip as shown in the figure. The thermal conductivity of sapphire

is similar to that of copper at low temperatures and it is also an electrical insulator.

The mass (and heat capacity) of the thermometer attachment system has been significantly

reduced which allows a faster time response than was possible before.

Another change in the A(H) system was made to allow better null temperature control

between the TEMPERING POST and the specimen. The condition T (specimen) - T (tempering

post) = 0 must be maintained in order to eliminate heat flow between the specimen and the

tempering post. A differential thermocouple is used as a sensor in the automatic control

circuit used to maintain this zero temperature difference. It is essential that the

thermocouple be electrically isolated from the system ground and yet be in very good

thermal contact with both the POST and the specimen. The thermal and electrical properties

of sapphire are again used to advantage as shown in figure 3. The previous method of

electrically isolating the junctions was to place cigarette paper between the wires and

the surfaces where the temperature was to be sensed.

The resistance of the CRT's had been measured using an a.c. lock-in amplifier as both

the bridge power supply and null detector. This system has been replaced, at least for

the present time, by a d.c. Wheatstone bridge. Resistance measurements made with the d.c.

system are more precise; however, current reversing procedures must be followed which

amounts to doubling the number of readings compared to the a.c. system.

2. 2 Materials

One alloy, S 304 00 (AISI 304), was tested during this reporting period. The composition

in weight percent of a specimen of this material taken from the same stock as the X(H)

specimen is as follows: C = 0.07, Cr = 17.82, Mn = 1.05, Ni = 8.94, P = 0.022, Si = 0.37,

S = 0.012, and the balance is Fe. The test specimen was annealed at 1000°C for 1.5 hours

in a vacuum of 0.133 Pa. The residual resistance ratio (RRR) for the specimen is 1.43.

The Rockwell hardness of material adjacent to the A(H) specimen is B93 in the unannealed

condition and B71 in the annealed condition.

The changes in the method of mounting the thermometers and in reading the thermometer

resistances appear to have increased the precision of the temperature measurements. Previously, the calibration of the CRT's by comparison to the calibrated germanium resistance

thermometer yielded an analytical representation with an imprecision of 5 mK for the CRT

nearest the TEMPERATURE CONTROLLED HEAT SINK and 9 mK for the CRT nearest the SPECIMEN

HEATER. The initial tests on S 304 00 indicate that these imprecisions have been reduced to

about 2 mK and 4 mK respectively.

The present data for the thermal conductivity of S 30400 at H = 0 and H = 6366 kA/m

(80 kOe) are shown in figure 4. Data were also taken at 796 kA/m (10 kOe) , but are indistinguishable from the H = 0 data. A comparison of the present zero field thermal

conductivity data with data for a similar material indicates that there are no significant

systematic errors present in the H = 0 determination. Table 1 shows the compositions of the

present specimen and a similar material for which there are low temperature thermal con- ductivity data [3].

3. Results and Discussion

i c /.
ELEMENT

S30400

COMPOSITION (wt. %)

S34700

Cr 17.82 17.17

Ni 8.94 11.52

Si 0.37 0.59

C 0.07 0.057

Mn 1.05 1.34

P 0.022 0.14

S 0.012 0.007

Ni + Ta 1.10

Fe bal. bal.

Table 1. Comparative compositions of two stainless steels.

A comparison of the thermal conductivities at H = 0 for these materials is given in table

2.

AS34700/XS30400

6 0.92

8 0.97

10 1.00

14 1.01

16 1.01

18 1.01

Table 2. Ratio of thermal

conductivities of two stainless steels at H = 0

.

Clark, et al. [4] found the RRR for their S 30400 specimen to be 1.42. The RRR for the

present specimen was determined to be 1.43. This excellent agreement indicates that the

specimen used by Clark and the present specimen are electrically similar.

The general effect of a magnetic field is to reduce the thermal conductivity as shown

in figure 4. The reduction from zero field values caused by a 6366 kA/m (80 kOe) field is

9.2% at 5.5 K and 6.7% at 19.75 K for S 30400. Similar data reported earlier [5] show the

corresponding reduction in conductivity for N07 718 to be 8% at 5.5 K and 3% at 19.5 K; for

S31000 the reduction is 20% at 5.25 K and 11% at 19.5 K. The field effect is seen to be

slightly greater in the S 304 00 than in the N07 718 , as would be expected. The Lorenz

numbers for the Ni-Cr alloys are typically higher than those for stainless steels. This

indicates a stronger lattice component of thermal conductivity and therefore a smaller

magnetic field effect. The present data exhibit a considerably smaller field dependence

of thermal conductivity than was observed for the S 31000 specimen. The variation in the

field effect on the two stainless specimens is not understood at the present time.

4. Conclusions

The zero field thermal conductivity data for S 30400 is very close to that taken for a similar steel in a high precision, zero field thermal conductivity system. The reduction

in X of S30400 caused by a 6366 kA/m magnetic field is 9.2% at 5.5 K and 6.7% at 19.75 K.
5o References

1. Sparks, L. L. and Fickett, F. R. , Materials research in support of superconducting

machinery, National Bureau of Standards, Boulder, Colorado, Cryogenics Division,

March, 1974 NBSIR-74-359, AD-780 596/3WM.

2. Sparks, L. L. and Fickett, F. R. , Materials research in support of superconducting

machinery, National Bureau of Standards, Boulder, Colorado, Cryogenics Division,

October, 1974 NBSIR-74-393, AD-A0004586.

3. Hust, J. G. , Thermal Conductivity, Electrical Resistivity, and Thermopower of Aerospace Alloys from 4 to 300 K: II AISI 347 Stainless Steel, unpublished NBS Report.

4. Clark, A. F., Childs, G. E. and Wallace, G. H. , Electrical resistivity of some engineering alloys at low temperatures, Cryogenics 10_, No. 4, 295 (August, 1970).

5. Sparks, L. L., Materials research in support of superconducting machinery, National

Bureau of Standards, Boulder, Colorado, Cryogenics Division, April, 1975 NBSIR-75-

810, AD-A012365.

List of Figures

Figure 1. Magnetothermal conductivity probe and magnet.

Figure 2. Modified method of attaching the carbon resistance thermometers

to the specimen.

Figure 3. Modified method of making thermal contact with the differential

thermocouple junctions.

Figure 4. Thermal conductivity of S 30400 stainless steel as a function of

temperature for H = 0 and 6366 kA/m (0 and 80 kOe respectively)

.
List of Figures

Figure 1. Magnetothermal conductivity probe and magnet.

Figure 2. Modified method of attaching the carbon resistance thermometers

to the specimen.

Figure 3. Modified method of making thermal contact with the differential

thermocouple junctions.

Figure 4. Thermal conductivity of S30400 stainless steel as a function

of temperature for H = 0 and" 6366 kA/m (0 and 80 kOe respectively).

157
158
-Heater Heater

Figure 3. Modified method of making thermal contact with

the differential thermocouple junctions.

^ 3

S30400

' r

n§

H = 0

H = 6366 kA/m (80 kOe)

J L 6 8 10 12 14

Temperature, K

16 18

Figure k. Thermal conductivity of S30400 stainless steel as a function

of temperature for H - 0 and 6366 kA/m (0 and 80 kOe respectively)
MAGNETOTHERMAL CONDUCTIVITY OF SELECTED PURE METALS AND ALLOYS

L. L. Sparks

Cryogenics Division

NBS - Institute for Basic Standards

Boulder, Colorado 80302

ABSTRACT

The magnetothermal conductivity program was initiated to determine

the effect of a magnetic field on the thermal resistance of technically

important metals. The experiments are done in magnetic fields up to

kA

6366 — (80 kOe) and cover the temperature range from 4 to 20 K. The

results of this study are presented for a Ni-Cr-Fe alloy, AISI 310 stainless steel, oxygen-free copper, and a high purity copper specimen. A

magnetic field typically increases the electronic thermal resistance and

thus lowers the total thermal conductivity of a metal. The magnetic field

effect at constant temperature is shown by the relative change in thermal

resistance, AW/W , where AW is defined as W , - W Our experimental

n— U HtF(J rl— U. data show that the effect of a 6366 kA/m (80 kOe) magnetic field on the NiCr-Fe alloy is to increase AW/Wu_ n by 0.08 at 5 K and 0.03 at 19 K. The

H—

U

corresponding values for AISI 310 are 0.28 at 5 K and 0.10 at 19 K. The

thermal resistivity of the better conductors was found to be affected more

strongly. A 6366 kA/m (80 kOe) magnetic field causes AW/W to increase

to 1.3 at 5.5 K and 0.8 at 19 K for oxygen-free copper; for the high purity

copper, AW/W increases to 3.4 at 5.5 K and 2.7 at 20 K. H— (J

Key words: Copper; low temperature; magnetothermal conductivity; nickel

alloy; stainless steel; thermal conductivity.

160
MAGNETOTHERMAL CONDUCTIVITY OF SELECTED PURE METALS AND ALLOYS*

L. L. Sparks

Cryogenics Division

NBS - Institute for Basic Standards

Boulder, Colorado 80302

INTRODUCTION

The objective is to determine the effect of magnetic fields on the

thermal conductivity of technically important metals. The need for this

information arises from the development of rotating machinery operating at

cryogenic temperatures. The existing literature on magnetothermal

conductivity, X(H), is concerned almost exclusively with scientific

materials, e.g., very pure materials and single crystals. A complete

bibliography of magnetothermal conductivity was given by Sparks and

Fickett'3"

The materials studied in this program are being used or have potential

use in superconducting motors and generators. Optimum design of these

machines, which must operate at low temperatures while in magnetic

fields, requires a detailed knowledge of how the thermal properties of

the constituent materials are affected by a magnetic field. The broad

material categories of interest include superconductor stabilizing

materials such as copper and aluminum, and structural materials such as nickel alloys, stainless steels, and metallic composites.

APPARATUS AND EXPERIMENTAL PROCEDURES

The experimental determination of the X(H) of metals in high fields is

complicated by the requirement that the specimen be contained in a region

of homogeneous field. For magnets of reasonable size, this restriction

necessitates small specimen lengths and resultant small temperature differences along the specimens. The X (H) probe employed in this study was designed to be used in a superconducting solenoid with a 3.8 cm bore and

a 2.5 cm homogeneous (1%) field sphere. Specimen lengths are therefore

limited to approximately 2.5 cm.

Supported by the Advanced Research Projects Agency (ARPA): Department

of Defense. ARPA Order No. 2569.

161
The principle components of the A(H) system are shown in figure 1. A

specimen is shown with its axis mounted parallel to the magnetic field.

This configuration is used to determine the longitudinal A(H), Provisions

for making transverse A(H) measurements (heat flow perpendicular to the

direction of the field) have been included in the system design.

The basic operation of the system involves balancing electrical power

supplied to three heaters with the heat leak to the liquid helium bath via

the THERMAL LINKS (capitalized parts refer to figure 1). The electrical

heaters were wound, one each, on the TEMPERATURE CONTROLLED HEAT SINK (TCHS)

,

the SPECIMEN, and the TEMPERING POST. The power supplied to the TCHS

determines the approximate temperature of the specimen; the SPECIMEN HEATER

is used to establish a temperature gradient along the specimen; and the

TEMPERING POST HEATER is used to reduce the temperature difference

between the specimen and the tempering post to less than + 5 mK. The

TEMPERING POST and TCHS heaters are automatically controlled during all

tests while a constant current is supplied to the SPECIMEN HEATER.

The thermometers used in the probe are three 1/8 watt, 100 ohm Allen

Bradley* carbon resistors (CRT's) and a single calibrated germanium resistor

(GRT). The CRT's are located, one each, in the TCHS, and the two THERMOMETER BLOCKS. The GRT is located in the TCHS and is used for zero- magnetic-field calibration of the CRT resistors. The effect of the

magnetic field on the resistance of the CRT's is taken to be that published

2 by Neuringer and Shapira . Resistance measurements on the two specimen

CRT's are made using a lock- in amplifier as both detector and power

supply for an AC bridge. Both absolute and difference resistance measurements are possible using this system.

The thermal conductivity of a specimen is computed from the specimen

geometry (Area/length), the specimen heater power (Q) , and the measured

temperature difference along the specimen (AT). The relationship of A to

these parameters is given by

X - A • (1)

The use of trade names of specific products is essential to a proper

understanding of the work presented. Their use in no way implies any

approval, endorsement, or recommendation by the National Bureau of

Standards.

162
A series of measurements of AT and Q at various fields* (0 < H < 6366 kA/m

or 80 kOe) and temperatures (4 < T < 20 K) result in the data presented in

this report.

MATERIALS

A total of four materials have been tested—Ni-Cr-Fe alloy UNS-N07 718 , oxygen-free copper UNS-C10200, stainless steel UNS-S31000, and a high purity

copper referred to as NBS-B stock 7. UNS-N07 718 was tested in the age-hardened

condition with a Rockwell hardness of C39 and an average grain diameter of

0.06 mm. The composition in weight percent for this material is: Ni = 54.57,

Cr = 18.06, Fe = 17.08, Nb + Ta - 5.12, Mo = 3.18, Ti = 0.85, Al = 0.44,

Mn = 0.29, Si = 0.24, Cu, C, and S < 0.1%; the residual resistance ratio

(RRR = R_,_ /R.„) is 1.06. The UNS-S31000 stainless steel was measured in 273K 4K

the cold worked condition; its Rockwell hardness is B77 and its RRR = 1.27.

The composition in weight percent for this material is: C = 0.13, Mn = 1.72,

P = 0.016, S = 0.018, Si = 0.84, Cr = 24.47, Ni = 20.44, and Fe balance. Both

polycrystalline copper specimens were vacuum annealed at 850°C for one hour.

RRR for the UNS-C 102 00 copper specimen is 107. The electrical resistivity

of the NBS-B stock 7 copper specimen was too low to be measured in the

3 A(H) system; however, extensive work by Fickett on the electrical resistivity

of this copper stock indicates that RRR = 1520 for heat treatments similar to

that of our specimen.

RESULTS

The effect of a magnetic field on the thermal conductivity of UNS-N07 718

is shown as a function of temperature in figure 2. Data are shown only for

0 and 6366 kA/m (80 kOe) fields; curves for H = 796, 1592, and 3183 kA/m

(10, 20, and 40 kOe respectively) fall between the two curves shown. The

estimated 10% uncertainty of a single data point would indicate that the

data at 0 and 6366 kA/m (80 kOe) are indistinguishable. However, the curves drawn through the data points indicate a definite trend to lower thermal

conductivities at 6366 kA/m (80 kOe) . The zero field data for UNS-N07718

4

is within 10% of the zero field data by Hust, et al. for a similar Ni-Cr-Fe

specimen.

Preliminary data for the thermal conductivity as a function of temperature for UNS-S 310 00 stainless steel are shown in figure 3. Again, only the

0 and 6366 kA/m (80 kOe) curves are shown.

* The International System of Units (SI) designation for magnetic field strength

is ampere per meter, and this unit is used throughout this paper. The more con- ventional unit of magnetic field strength, the oersted, is also given. Conversion

of oersteds to amperes/meter is accomplished by multiplying oersteds by 1000/4tt.
The thermal conductivity and magnetic field effect are shown in figure

4 for UNS-C102 00 copper. To my knowledge there are two other sources of zero field thermal conductivity data for similar copper in the literature.

Powers, et al."* give no RRR for their specimen, which has a conductivity

about 50% lower than these data at 20 K. Hust and Giarratano find a conductivity about 90% higher than these data for a specimen whose RRR = 230. This

range of results is indicative of the sensitivity of the thermal conductivity in this temperature range to trace impurities and heat treatment.

Fevrier and Morize^ measured the magnetothermal conductivity of two copper

wires with RRR = 62 and 162 near 4 K. Our data fall generally between

the curves for their two specimens as one would expect.

Data for NBS-B stock 7 copper are shown in figure 5. The experimental

points at H = 0 represent averages while the experimental points for H

^ 0 are single data points. The extremely high conductivity at zero field is near the limit of what can be measured in the present system.

Repeat runs and data averaging were used to increase the reliability of

the zero field results. The source of the curve drawn through the H = 0

data will be discussed later in this paper.

DISCUSSION

Our data indicate that a 6366 kA/m (80 kOe) magnetic field reduces the

thermal conductivity of our UNS-N07 718 specimen by 8% at 5.5 K and by 3% at

19.5 K. An estimate of the electronic contribution to the thermal conductivity

of this alloy can be made using the Lorenz ratio (L = pA/T) given by Hust

8 —8 2 2 and Sparks, and L

q = 2.4 x 10 V /K . This estimate indicates that the

electrons carry roughly 25% of the heat flow at 5 K and 18% at 19 K. The

Lorenz numbers for stainless steels are typically lower than those for NiCr-Fe alloys by approximately a factor of 2. This would suggest a larger

electronic component in stainless steels, and consequently a larger field

effect. The effect of a 6366 kA/m (80 kOe) field on the UNS-S31000 stainless

steel specimen is to reduce the conductivity by 20% at 5.25 K and 11% at

19.5 K as determined at the extremes of temperature as shown in figure 3.

At intermediate temperatures the field appears to have little effect on the

thermal conductivity, i.e., the H = 0 and H = 6366 kA/m (80 kOe) data are

the same within the stated experimental uncertainty. This behavior was un- expected and is unexplained at the present time. Further work on similar alloys

should help to determine whether the observed behavior is due to some scattering

phenomenon or possible experimental error.

164
The zero field thermal conductivity of the present data (figure 4)

for UNS-C102 00 copper exhibits a nearly linear decrease with decreasing

temperature. This is the behavior expected for impurity scattering of

electrons. The effect of the magnetic field is to reduce the thermal

conductivity of our UNS-C102 00 copper specimen by 45% at 6366 kA/m (80 kOe)

and 19 K and by 60% at 6366 kA/m (80 kOe) and 5 K.

The very high thermal conductivity of the NBS-B stock 7 copper specimen

causes experimental difficulties that result in lower accuracy in the zero field runs than was the case for the lower conductivity materials. The

9 solid curve shown in figure 5 for H = 0 was derived using the relationship

1/A = W = AT

n + B/T = 3.02 x 10~ 8

T

2,55 + 4.7 x 10~4

/T. (2)

This equation represents the experimental data within the experimental accuracy

the units of W are mK/W. The effect of a magnetic field is included in figure

5. The anomalous decrease shown in this figure for A at 1592 kA/m (20 kOe)

and high temperatures is probably not real; the single data point at 21 K

appears to be bad, but no reason can be found to discard it. Figure 6

presents the relative change in thermal resistivity

AK/K

H-0

= (A H=0 " V/A

H

(3)

as a function of magnetic field. The high temperature, the 1592 kA/m (20 kOe)

data have been ignored in this figure due to its anomalously low value.

The result is that, although the exact placement and shape of the knee is

somewhat questionable, there is no doubt a rapid decrease in slope of AW^_^

in the 796 to 3183 kA/m (10 to 40 kOe) range. A similar plot of AW/WU_. n—

U

for UNS-C102 00 copper is shown in figure 7; there is no rapid change in the

slope of AW/W _ for this material. A magnetic field should affect the

H—

U

electronic conductivity strongly whenever oot >_ 1 where 03 is the angular

frequency and T is the relaxation time. oj is proportional to the magnetic

field as indicated by the cyclotron relationship CO = eH/mc. Further,

T (high purity copper) > T (low purity copper), so that with increasing

fields ojt becomes large for the high purity stock 7 copper before it does

for the UNS-C 102 00 copper specimen; consequently, the effect of a magnetic

field should be seen at lower fields for the stock 7 copper than for UNSC 1020 0 copper. A comparison of the data in figures 6 and 7 shows this to be

the case.

165
In conclusion, the effect of a magnetic field on structural alloys,

such as the Ni-Cr-Fe and stainless steel reported here, is relatively small.

The results are not inconsistent with estimates of the possible field

effects using Lorenz ratio data. Good conductors such as the two copper

specimens show greater field effects which are to be expected due to the

large electronic component of the thermal conductivity.

REFERENCES

1. L. L. Sparks and F. R. Fickett in: "Materials Research in Support of

Superconducting Machinery," Nat. Bur. Stand. NBSIR-74-359 (March 1974).

Available from NTIS, AD-780 596/3WM.

2. L. J. Neuringer and Y. Shapira, Rev. Sci. Instrum. , 40(10): 1314 (1969).

3. F. R. Fickett, "A Preliminary Investigation of the Behavior of High

Purity Copper in High Magnetic Fields," Annual Report to INCRA

(August 1973).

4. J. G. Hust, D. H. Weitzel and R. L. Powell, J. Res. Nat. Bur. Stand.

(U.S.), 75A(6): 269 (1971).

5. R. W. Powers, D. Schwartz and H. L. Johnston, "The Thermal Conductivity

of Metals and Alloys at Low Temperatures I. Apparatus for Measurements

between 25° and 300°K. Data on Pure Aluminum, OFHC Copper, and L Nickel,

Ohio State University, TR 264-5 (1951).

6. J. G. Hust and P. J. Giarratano in "Materials Research in Support of

Superconducting Machinery," Nat. Bur. Stand., NBSIR-74-393 (October 1974)

7. A. Fevrier and D. Morize, Cryogenics, 13: 603 (1973).

8. J. G. Hust and L. L. Sparks, Nat. Bur. Stand. (U.S.) Tech. Note 634,

(February 1973)

.

9. H. M. Rosenberg, Low Temperature Solid State Physics, Oxford University

Press, London (1963), p. 117.

]66
I 76X 1969 I

Fig. 1. Magnetothermal conductivity probe and magnet.

167
5.0 i r i—i—

r

4.0

3.0

UNS-N07718

H = 0

2.0

1.0

H = 6366 kA/m

(80 kOe)

10 12 U 16 18 20 22

TEMPERATURE , K

Fig. 2. Thermal conductivity of the Ni-Cr-Fe alloy UNS-N07 718 as a function of temperature for H = 0 and 6366 kA/m (80 kOe)

.

TEMPERATURE , K

Fig. 3 Thermal conductivity of UNS-S 31000 stainless steel as a function

of temperature for H = 0 and 6366 kA/m (80 kOe).
rig. 5 Tbexnal conductivity of "nigh parley copper (BBS » 1320) as a

foactioct of tenperatare for E = 0, 796, 1592, 3183, and 6366 sJJtl (0, 10, 20, 40, and SO kOe respectively)

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NBSIR

SEMI-ANNUAL REPORT ON MATERIALS RESEARCH

IN SUPPORT OF SUPERCONDUCTING MACHINERY

THERMAL CONDUCTIVITY

J. G. Hust

Cryogenics Division

Institute for Basic Standards

National Bureau of Standards

Boulder, Colorado 80302

October 1975
Summary: Thermal Conductivity

Thermal conductivity data are reported for boron-epoxy composites. Comparisons are made with existing literature data and material variability is discussed.
THERMAL CONDUCTIVITY

J. G. Hust

Cryogenics Division

Institute for Basic Standards

National Bureau of Standards

Boulder, Colorado 80302

1 . Intr oduction

At the request of Mike Jones for Watervliet Arsenal and in conjunction with the

ARPA-NBS superconducting machinery program, the thermal conductivities of three

specimens of boron-epoxy composite were measured. Both a transverse and a longitudinal

specimen were measured using a fixed-point apparatus and a single longitudinal

specimen was measured using a variable-temperature apparatus.

2 . Apparatu s and Spe cimen Characterization

A fixed-point thermal conductivity apparatus described by Hust [1,2] was used to measure transverse and longitudinal boron-epoxy composite specimens. The specimens

were 2.54 cm in diameter and 3.56 cm in length. Thermal conductivity measurements

with this apparatus are performed near the boiling point of helium (4 K) and nitrogen

(76 K) , the sublimation point of CO2 (194 K) , and the ice-point (273 K) . The uncertainty of these measurements is estimated to be + 10%.

The specimens were fabricated from 5.6 mil boron-epoxy pre-impregnated tape. The

boron volume fraction was about 48%. The void fraction is stated to be less than

2% by the supplier. A cursory examination of the material suggests that the void

fraction is appreciably less than that.

A longitudinal specimen, with the same characteristics as above, was also measured in a variable-temperature apparatus. The specimen was 23 cm long and rectangular

in cross-section (1.74 cm x 2.56 cm). The variable-temperature apparatus is designed

to obtain data as a function of temperature from 4 to 300 K and is more accurate than

the fixed point apparatus. This apparatus is described in detail by Hust et al. [3].

3. Results

.'ieasurements were performed on a longitudinal specimen (i.e., in the direction of

the boron fibers) of boron-epoxy composite at temperatures from 4 to 300 K using the

var Lable-temperature apparatus. The experimental data are listed in table 1 and

i ] Lustrated in figure 1. The following equation was least-squares fitted to these

data using orthonormal fitting techniques:

r i+1

£n>, = y a.[£nT] O)

i=l 1

is thermal conductivity in W*m ^"*K ^ and T is temperature in kelvin. The deviations

the experimental data from this representation are listed in table 2 and are illustrated in figure 2. The parameters obtained by least-squares fitting are listed in

tabic 3. Tabular values of A as computed from the above equation are listed in

table 4 and illustrated in figure 3.

173
Measurements were conducted on both a longitudinal and a transverse specimen using

the fixed-point apparatus. The resulting values of thermal conductivity are listed

in table 5 and plotted in figure 3. The uncertainty of the smoothed data obtained with

the variable-temperature apparatus is estimated at + 4% and the uncertainty of the

fixed-point apparatus data is estimated at + 10%.

Literature data for similar composites have been compiled by Kasen [4] . These

data [5,6] are included schematically ii! figure 3 for comparison to the data from this

research. Reasonable agreement is observed between the literature data and the

longitudinal specimen measured in the variable temperature apparatus and the transverse

specimen. However, the longitudinal specimen measured in the fixed-point apparatus has

a conductivity lower by about 50% than the other longitudinal specimen.

At ambient temperature the thermal conductivity of crystalline boron is 20 W*m ^"K ' and the thermal conductivity of epoxy is typically about 0.2 W'm- -'- •K

- -'- . Since these

composites are about 50% boron one would expect a conductivity near 10 W*m- ^ •K

- -'- instead of 1 - 2 W*m~l»K-"'" . The observed results suggest the presence of strong phonon

scattering in the boron fibers. These fibers are fabricated by vapor deposition of

boron on a 0.5 mil diameter tungsten wire substrate. (See reference 7 for details).

During the fabrication process compounds of tungsten and boron are formed at the center

of the final filament and the major outer portion is reported to be pure amorphous or microcrystalline boron. The center portion has a lower thermal conductivity, but since

it represents only a few percent of the cross section of the total filament, its effect

on reducing conductivity can be neglected. If the outer boron portion of the filament

is in a highly disordered or amorphous state, the thermal conductivity would be con- siderably reduced from that of crystalline boron. It is suggested in reference 7 that

the boron is amorphous or microcrystalline structure with an effective crystal size

O of near 20 A. Boundary scattering of the heat carriers (phonons) from such small

crystals is sufficient to reduce the thermal conductivity considerably since the phononphonon scattering mean free path at ambient temperature is near 25 A. The phonon-phonon

mean free path, t, is estimated from bulk boron thermal data and the classical equation

^ = 1/3 Cv£. It is likely that significant crystal size variations occur between

different production runs of boron filaments. Thus, one would expect large batch-tobatch variations in thermal conductivity of the longitudinal specimens, which is con- sistent with the observed results.

In the transverse direction one can estimate the conductivity of the composite

from a knowledge of the epoxy conductivity and the relative amount of epoxy present.

From such considerations for this composite one obtains a value of conductivity of

about 3 to 4 times greater than the conductivity of epoxy. This corresponds to 0.6

to 0.8 W , m~l , K

- -'- at ambient temperature which is in reasonable agreement with the measured value of 0.6 W*m~l , K

-

-'-. Similar agreement is obtained at low temperature.

174
References

1. J. G. Hust, Low temperature thermal conductivity measurements on longitudinal and

transverse sections of a superconducting coil, Cryogenics 1_5, No. 1, 8-11 (1975).

2. J. G. Hust, Thermal conductivity fixed-point apparatus, NBS Laboratory Note 74-3,

9 pp. (1974).

3. J. G. Hust, R. L. Powell and D. H. Weitzel, Thermal conductivity standard reference

materials from 4 to 300 K. I. Armco Iron: Including apparatus description and error analysis, J. Res. Nat. Bur. Stand. 7Uk, No. 5, .673-90 (1970).

4. M. B. Kasen, "Properties of filamentary-reinforced composites at cryogenic temperatures", Composite Reliability, ASTM STP 580 pp. 586-611 (1975).

5. J. Hertz, J. L. Christian and M. Varlas, "Advanced composite applications for

spacecraft and missiles, Phase I final report, Volume II: Material development",

AFML-TR 71-186, Vol. 2, March 1972 (AD 893 715L)

.

6. J. P. Gille, "Development of advanced materials for integrated tank insulation

systems for long-term storage of cryogens in space", NASA CR-102 570 (Final) Sept.

1969 (N 70-23348) .

7. Modern Composite Materials, edited by L. J. Broutman and R. H. Krock, AddisonWesley Publishing Company, Reading, Mass. (1967).

175
NOTES RELATING TO TABLES

THfc DATA

THE COMPUTER

FOLLOWING IS

LISTED AFT, in PART, CuRD IMAGES OF FOR DATA PROCESSING. ThESE DATA ARE

A LIME BY LINE EXPLANATION OF THESE

EXPERIMENTAL DATA AS READ INTO

GENERALLY UNLA3ELLEO. THE

TABLES.

TABLE 1

LINE 1-

LINE 2-

LINE 3-

LINE

(GRADIENT DATA)

DATA ITENTIFICAT1 ON EMFS OF THERMOCOUPLES 1 THRU 6 ( MICRO VOLT S

)

PLATINUM RESISTANCE THERMOMETER CURRENT (MILLIAMPS), REFERENCE

PLATINUM TH r chO'ETER VOLTAGE (MICROVOLTS), FLOATING SINK PLATINJM

THERMOMETER VOLTAGE ( '-'IbRGVOLTS) , GERMANIUM T HE t\M OMc T cR CURRENT

(MICROAMPS), REFERENCE GERMANIUM THERMOMETER VOLTAGE (MICROVOLTS),

FLOATING SINK GERMANIUM THERMOMETER (MICROVOLTS)

SEEBECK FMF (MICROVOLTS), SPECIMEN VO LT AG r ( M IC "0 VOLTS) , SPECIMEN

CURRENT (MILLIAMPS), SPECIMEN HtATER VOLTAGE (VOLTS), SPECIMEN

HEATEP CURRENT (MILLIAMPS), CRYOGENIC 3 AT H PRESSURE ( Nfc OF HG ) , ROOM TEMPERATURE (DEC-EES C), CRYOGtNIC BATH CODE (1=LIUUI0 HELIUM,

2=LIQUID HYDROGEN, 3=lIQUI0 NITROGEN, i» = DRY ICE-ALCOHOL, C =IC; dATER)

TAeL:

THlSE DATA ARE StM I- PROCESSED COMPUTER OUTPUT. TEMPERATURE IS KELVIN,

THERMAL CONDUCTIVITY IS IN WATTS PER MITER PER KELVIN, ELECTRICAL

RESISTANCE IS IN OHMS, AND TH e PMO VOLT AGE IS IN MICROVOLTS.

BORON-EPOXY (LI)

NUMBER OF CALIBRATION, I SO THERMAL , A NO GRADIENT RUNS, mN ) THERMOCOUPLES

6o -C 11 6

THERMOCOUPLE POSITIONS ( 0 M

)

O.OCQO 2.5*GC 5 . 0 80 C 7.620G 10.16CO 12.rCuC 15. 2^00 17.7800

SPECIMEN OIAMETERS BETWEEN THE RMOCUUPLf S (CM)

2.3P1E 2.3815 2.3815 2.381? 2.3815 2.3315 2.3815

176
TA8LE 1. BASIC SEN I-PROCESSEO TEMPERATURE GRADIENT OAT A FOR

BOPON-tPOXY ti.1)

THERMAL CONDUCTIVITY DATA FOR B-EPOXY(Ll) 20 * A Y 75 1545

11.61+

2.0005

-0 .o-e

PR T 0.000

20 .80

34.8 0 -ft. CO

REFERENCE TEMP

G»T

4. 158

28. 7t) 35. 79

25.78 9.9880

-0.O-8 .6740

FLOATING SINK TE^P

PPT CRT

Q . 00 C h. o52

=.1.99 48.22

9657.50 7598.50

3.3230 632.7

SPECIMEN

RESISTANCE

9 .

11

>3 .7;

24. f) HEATER

POWE c ..22431E-G2

53.72

1. 0

THERMAL CONDUCTIVITY DATA FOK F-'_P0XYIL1I 20 MAY 75 1620

45.0 3 66.97

•0 .0 000 -0 .0 0 -0.90 -0.0C

REFERENCE TEMP

PR T GRT

0 . 0 30 -.194

85 .00

-0.00

101.14

9.9o3C

-0.03 1.356C

FLOATING SINK TEMP

PPT GRT

C.GOG 6 . 172

115 .12

9565.50

123.43

4<+2 :i . 8 0 6.6972

SPECIMEN

RESISTANCE

9 .

632. 7

12 140 .57

24.0

hz. ATER

POWEF

90 B14E-02

151.52

1.0

THERMAL CONDUCTIVITY DATA FOR P-EPOXY(Ll) 20MAY 75 1715

127 .92

2.0C1Q

-0 .00

PR T 0. 000

167 .36

34.70

-0.00

REFERENCE TEMP

GRT

4. 289

2CQ.18 230.12

t-0.16 9 . 9830

-0 . 00 2 .1.880 FLOATING SINK TEMP

PR T GRT

G.GOG 1C. 31

3

256.23 230.38

9113.01 15L9.5G

12.2823 635.

u

SPECIMEN

RE il ST ANCE

0 .

13

3 03 . 32

HEATER

POWE P •3G553E-G1

32-. 92

1. 3

THE R MA L CONDUCTIVITY DATA FOF ?-EPOXY(Ll> 2 3 M A Y 75 17*0

265 . 04

2.0C10

-0 .00

PEFEP'NC:

PR T 0. COO

340 .28

34.66

-O.OO

TF^P

GRT

4 . *»6 0

403 . 64

10>.94

-0 .00

FLOATING

PR T lb. 3 C 7

463 .67

9.9830

4 . 5 14 C

' SINK TEMP

GRT

15. 931

517 .90 569 .72

14

6 2 3 • o o 6470 . 0 0 65- .80

22.2553 635. C 2*,

SPECIMEN HEATE 5 RESISTANCE POWE 5

6c9 . 3 5

1. C

0 . . 1 3 C-oE O J THERMAL CONDUCTIVITY C AT A FOK F-EPOXr(Ll) 20MAY 75 1945

698.97 751.82 802.25 85b. 47

2 . 0 00 3 -0 . 0 0 274*3. 50 9 . 9880

-0.00 -0.00 -0.0C 4.6560

REFEPENCE T FMP FLOATING SINK TEMP

PR T GRT PRT GRT

0.099 4.810 43. 339 0 . 000

907.42 • 953 .15

7350.33 136.18

22.928J 635.0

SPECIMEN

RESISTANCE

0 .

15 1010 .90

24. G HE A TE r POWE R

. 10 <o7jt. *3 u

1062.44

1. b

THERMAL CONDUCTIVITY DATA FOR d-tPOXYCLl) 17 MAY 75 2215

113.31 163.22

2.0G20 9287. DO -0.0G -0.00

REFEPENCE TEMP

PR T G<?T 75. 90 5 C. 00 0

211. 2h

1 0280 .00

-0 .00

FLOATING

°R T 80. 332

262.72

-C . G 00 0 4. 3020

SINK TEMP

GRT

0 . GOO

312.72

-0 .00

21.2100

361 .91

-0.00

63 3.5

SPECIMEN

RESISTANCE

0 .

-13 . lc

25.3

HEATED

POWE R

. 91245E-01

>1. 68

3. 3

177
TABLE 1. (CONTINUED)

THERMAL CONDUCTIVITY DATA FOR B-EPOXY(Ll) 20MAY 75 1040 10 T296.88 14b4.3*5 IF2¥n.TJ 179/. 40 1963.96 2T24.4 0 22 93. 2E 2 ¥5 0 . 64

2.0010 9318.00 23035.00 -0.0000 -0.00 -0.00

-0 .00 -0.00 -0 .0 0 7.274 0 35.7000 634.5 23. 8 3.

0

REFERENCE TEMP FLOATING SINK TEMP SPECIMEN HEATER

prt -grt" pk r cm— —RTsrsrancF porer

76.067 0.000 138.550 0.000 0. .25968E + 00

THERMAL CONDUCTIVITY DATA FOR B-EPOXY(Ll) 18 MAY 751155 8 34B.66 468. 5* 584.84 712.57 835.65 956.94 1083.82 1203.9*

2 . Q 00 0 9290 .20 1 2589 .00 -0.0000 -0 . 00 -0 . 00

-O .00 -0.00 -

^U.O0 6.5260 32.0900 633.7 Z5.5 3.

0

REFERENCE TEMP FLOATING SINK TEMP SPECIMEN HEATER

PRT T5RT PRT GRT RESISTANCE POWER

75.962 0.000 90.790 0.000 0. .20942E+00

THERMAL CONDUCTIVITY DATA FOR B-EPOXY(Ll) 14 MAY 75 1345 3 rZSV5~0 I95T8D ?B3r20 33TJ.8TJ 403.30 469.70 537TT0 6TrO.?0

2.0025 34565.00 35520.00 -0.0000 -0.00 -0.00

-0.00 -0.00 -0.00 4.5540 22.3600 -0.0 23.0 4.0

REFERENCE TEMP FLOATING SINK TEMP SPECIMEN HEATER

PRT GRT PRT GRT RESISTANCE POWER

193.251 0. 000 197. 391 0. 000 0 . .10183E*00

THERMAL CONDUCTIVITY OATA FOR B-tPOXY(Ll) 15 MAY 75 2140 5 591.89 776.53 956.91 1149.69 1335.47 1511.59 1692.39 1859.22

2.0010 34720.00 39595.00 -0.0000 -0.00 -0.00

-0.00 -0.00 -0.00 7.4000 36.3000 -0.0 23.0 4.0

REFERENCE TEMP FLOATING SINK TEMP SPECIMEN HEATER

PRT "GRT PRT"— GRT RF SI STANCE PTJWER 194. 118 0. 000 217. 087 0. 000 0 . .26862E+00

THERMAL CONDUCTIVITY DATA FOR B-EPOXY(Ll) 7 MAY 75 835 1 48.78 102.73 153.90 208.29 260.50 310.31 360.70 407.00

2.0000 51010.00 51420.00 -0.0000 -0.00 -0.00

-O.00 -0.00 _ -0 .00 3.9TJF8 "18.9~45iJ~ -0.0 ~-O.0 5.0

REFERENCE TEMP FLOATING SINK TEMP SPECIMEN HEATER

PRT GRT PRT GRT RESISTANCE POWER

178
TABLE 2 THERMAL CONDUCTIVITY DEVIATIONS FOR

BO RON- EX T < L1)

T H ER M AL COND UCTIVITY D ATA FO R B-EP OXY tLl) 2 0H A Y 7 5 1 5 *»5 11 MEAN TEMPERATURE OBSERVED CALCULATED PERCENT

TE MP ERATURE DIFFERENCE " TH E RMAL —

T

HERMAL —DEVI AT 10

N

CONDUCTIVITY CONDUCTIVITY

5. 522 ~ ,678 .18 9E »0 0 .19 1 E » 00 -1.1 "

6. 133 .5<*6 .23<*E«-00 .230E+Q0 1.9

6.E>5t» .V35 ,r5^£*tr0 ,265E*00 -2,7

7.112 .'21 .30<*E*Ot) .297E*00 2.3

. . 7. 518 .390 .328E*00 .325E»00 .9

7.898 .370 .3<*5E*00 .352E*00 -1.9

8.2<* 3 r3Et .<+00E*00— - . 376E*-00 6TU

THERMAL CONDUCTIVITY DATA FOR D-EP~OXYT LI) 20MAY 75 1620 TE

MEAN TEMPERATURE OBSERVED CALCULATED PERCENT

TEMPERATURE DIFFERENCE THERMAL. THERMAL OEVIATION

CONDUCTIVITY CONDUCTIVITY

8. 2» 0 HIT .36 5 E» 00 .

3

76E »0 0 -2". 9

9.510 1.122 .<*61E*00 .t*66E*-00 -1.0

10.576 1.009 .5r3E*00 .5t*tE*0U -E . 5

11. 506 . 853 .607E+00 . 607E*00 .1

12.319 .773 .6T0E + 0O .663E + OF 1.0

13.071 .731 .709E+-0Q .715E+Q0 -.9 - 1r3. 75 8 r€r«n» .8 0 5E»0 0 " . 762E» 0 0 ~ 5.3

THERMAL CONDUCTIVITY DAT A FOR B-EPOXYfLU 20 MA Y

" 75" 1715 15

MEAN TEMPERATURE OBSERVED CALCULATED PERCENT

TEHPERATtTRE DIFFERENCE THERMAL "THERMAL OEVIATION

CONDUCTIVITY CONDUCTIVITY

- 13. 886— - 2. 3 22 .750E » 00 " .77 IE* 00 _ -2.7

15.99W 1. 893 .920E«-&0 .912E+0Q .9

17. 812 1.71*3 .999E»-00 . 103E*01 -2.9

19. 3 1.52G .115£*-01 .113E + 01 1.5

20.90<* l.i»0 1 .12«*E*01 .122E*01 2.2

22. 276 1.3<.i* .130E+01 . 129E«-Q1 .3

~ 23. 56** ~ 1. 232 ' .1«*1E*01 —'. 13 6E*01 3.8

179
TABLE 2 (CONTINUED)

THERMAL CONDUCTIVITY n AT A FOf B-EPOXY(Ll) 20 MA Y 75 1740

MEAN TEMPERATURE OBSERVED CALCULATED PERCENT

TEMPERATURE DIFFERENCE THERMAL T HERM AL DEVIATION

CONDUCTIVITY CONDUCTIVITY

22. 996 <+. 390 .130E+01 . 13 3E*01 -2.1

27. 053 3.73C .153E+01 . 153E+-01 . h 33. 71^ 3. 576 .160E* 01 . 16 7E* 01 -<*.5 Zk, 132 3. 257 .1 76E + 01 . 178E«-01' -1.2

37. 30^ 3.08b .1 86E* 01 . 186E*01 -.2

UQ. Z7k 3. 051 .188E+Q1 . 192E*01 -2.1

»*3. 3<*9 2. 90 C .198E+ai . 196E«-C1 . 3 THERMAL CONDUCTIVITY DATA FOR B-EPOXY(Ll) 20MAY 75 1945

MEAN TEMPERATURE OBSERVED CALCULATED PERCENT

TEMPER ATURE OIFFLRENCE THERMAL THERMAL DEVI AT IO N CONDUCTIVITY CONDUCTIVITY

<*7. 792 3. 007 .202E+01 . 20 0E+01 1.1

50. 825 3.058 .1 99E*01 . 20 2E4-01 -1.5

53. 739 2.771 .220E+ 01 . 20 3E + 01 7.6

56. 667 3. ORE .197c* 01 . 20 4tVrii -3.2

59. 783 3. 1C6 .193E+01 . 204E+01 -5.3

62. 793 2. 87 3 .212E+01 . 20 3E*01 k . G 65. 638 2. 918 .209E* 01 . 20 3E>01 2. 7 THERMAL CONDUCTIVITY DATA FOR b-EPOXY(Ll) 17 MAY 75 2215

MEAN TEMPERATURE OBSERVED CALCULATED PERCENT

TEMPER ATURE DIFFERENCE THERMAL THERMAL DEVIATION

CONDUCTIVITY CONDUCTIVITY

83. <*69 2. 70 2 .193E*01 . 193E*01 -2.9

86. 118 2. 595 .200E4-01 . 197E4-D1 1.5

88. 785 2. 739 .190E4-01 . 19 7Ef01 -3.6

91. ^89 2. 66 7 .195E+01 . 196E + 01 _ c

. 9<*. 126 2. 610 .199E* 01 . 195E+01 2. 0 96. 777 2. 591 .193E«-Q1 . 195E + 01 -.8

99. 393 2. 5<+2 .205E+01 . 19-fEi-Ol 5. 1

180
TA8LF 2 (CONTINUED)

THERMAL CONDUCTIVITY DATA FOP B-EPOXY(Ll) ^ n m a v 7C 1 n i, n cUnUT f z> IUhU

MEAN TEMPER AT URF OBSERVED CALCULATED PERCENT

TEMPERATURE DIFFERENCE THERMAL THERMAL DEVI ATION

CONDUCTIVITY rnwnnrT tu ttv

lh7. 398 8.133 .182E* 01 1flQhi.fi! ** • J 155. 381 7. 833 .1B9E+ 01 1 b op a. n 1 i 19 * U 1 • 1 163. 386 8.176 .laiEf oi i 10 ?L t U 1 H 1 171. Hh5 7. 9^2 .186E«-01 • 10 ot U X

— a 179. 281 7.731 • 192E+01 1 ft rtr 4 n 1 • 1 0 Ol tui c • u 187. 100 7. 907 .187E* 01 1 R 4-0 1 n 19<*. 710 7.313 .202E+ 01 • 10( 7 A

THERMAL CONDUCTIVITY C AT A FOR Q-EPOXY ( LI

)

1 Q M A V 7 C 1 1 C C

1 0 "AT f JllJ>

MEAN TEMPERATURE OBSERVED CALCULATED PERCENT

TEMPERATURE DIFFERtNCE THEPMAL T HERMAL OE VIATION

CONOUCTI VITY CONDUCTIVITY

97. 998 6. 300 ,190E*01 . 19 5E ->01 -2.6

10<*. 18S 6. 07« •196E*01 . 19 5E«-01 1.6

110. U88 6. 522 .183E+ 01 . 192E + 01

116. 885 6. 272 .190E*Q1 . 19 IE +01 - . D 123. 090 6. 139 .195E+01 . 191- +01 1. 9 129. 32U 6.330 .1 89E*01 . 190E «-01 -.9

135. hot* 5.950 .201E*01 . 193E*C1 5.<« THERMAL CONDUCTIVITY DATA FOR B-EPOXY(Ll) ih MAY 75 13f+5 MEAN TEMPERATURE OBSERVtO CALCUL ATEO PERCENT

TEMPERATURE DIFFERENCE THERMAL THERMAL UEVI AT ION

( CONOUC TI VITY CONDUCT IV IT Y 200. 766 3. 20 6 .181E+01 . 136E+01 -2.9

203.936 3. 13U -.lB5c*fll . 186£*01 -."5 207. 16W 3. 321 .1 75E+01 . 186E+01

210. <4l2 3. 175 .1 83£«- CI . 136E+01 -1.6

21 3. 525 3.05 0 . 190E* 01 . 18&E+01 2. 5 216. 593 3. 086 .1 38L+ 01 . 18 5F. fOl 1. 3 219. 583 2.895 .201E*01 . 185E+01 7. 7

10

181
TABLE 2, (CONTINUED)

THERMAL CONDUCTIVITY DATA FOfc B-EPOXY(Ll) 15 MAY 75 2140

MEAN TEMPERATURE OBSERVED CALCULATED PERCENT

TEMPERATURE DIFFERENCE THERMAL THERMAL DEVIATION

CONDUCTIVITY CONDUCTIVITY

225. 745 8. 433 .182E*01 . 18 5Ef 01 -1.8

. Z34. 10h 8. 285 .iB5E>0i . 185E+01 . 2 242. 656 8. 819 .174E+G1 . 18<»E*0i -6.2

251. 275 8. 418 .182E+01 . 18<+E«-01 -1.3

259. 512 8.056 .190E+U1 . 185E+01 2.9

267. 609 8. 139 .188E+01 . 185E*Q1 1.5

275. 392 7. 40 7 .207E+01 . 186E+01 1C. 0 THERMAL CONDUCTIVITY DATA FOR B-EPOXY(Ll) 7 HAY 75 835

MEAN TEMPERATURE OBSERVED CALCULATED PERCENT

TEMPER ATURE DIFFERENCE THERMAL THERMAL DEVIATION

CONDUCTI VITY CONDUCTIVITY

276. 682 2. 4i«* .175£>Q1 . 186E+01 -o.7

279. 237 2.29 5 .184E4-01 . 187E+01 -1.6

281. 612 2. <*55 .172t* 01 . i87Ef01 -8.9

284. 005 2. 332 . 181E+ 01 . 18 8E + 01 -3.7

286. 287 2. 232 .189E+01 . 183t»0i . 5 288. 533 2. 260 .187E+01 . 189E4-01 -1.0

290. 687 2.0*9 .2 Q6E* 01 . 189Ef 01 8.1

TABLE 3. Parameters in equation 1 for a longitudinal

boron-epoxy composite specimen

i a

.

l

1 -17.065713

2 29.019209

3 -21.740071

4 9.1816220

5 - 2.3311507

6 0.35284329

7 - 0.029346836

8 0.001032607

182
TABLE 4. Thermal conductivity values of a lonj;itudinal

boron—epoxy composite specimen as calculated from equation 1

T1 A T A (K)

i -l (W'm *K ) (K) (W m

_1

'K

o n ill u . ZZ± 60 2 . 04

7 0.289 70 2.02

8 0.359 80 1.99

9 0.430 90 1.96

10 0.501 100 1.94

n a/, i Ton 1ZU 1 . 91

14 0.779 140 1. 90

16 0.912 160 1.89

18 1.04 180 1.88

20 1.16 200 1.86

25 1.43 220 1.85

30 1.65 240 1.84

35 1.80 260 1.85

40 1.91 280 1.87

45 1. 98

50 2.02

TABLE 5. Thermal conductivity values of a transverse and a longitudinal boron-epoxy composite specimen as measured

using the fixed-point apparatus

T X

(K) (Wm"1 ^"1

)

7.83 0.178

82.0 0.913 . . ,. . _ in, / Longituamal Specimen

iy i . ii

280 1.02

14.7 0.174

) 86.7 0.466 I _ 2qq q

J Transverse Specimen

280 0.581 |

183
4.8

2.8"

* 1.0

~ .6 a

O

a:UJ

.2"

I i I > i t i « i %

t

28

i > i i » i » i i I 48 SO 18 108 TEMPERATURE (Kt 1 v I

n

288 488

Figure 1 - Experimental data on longitudinal boron-epoxy composite specimen

184
» I • f -I » •

4

-fLxJ or

or

UJ

a.3T

co

•rTM 4J

(T3 C7

CUE

O

s_ 4-

fO -M

fO "Or—

ro

+->

CDE• r—SOlCLX

at 4-O

too

•r—

•r— cu > +-) cu XI too

>> CLE

•r" o

> u

pTM

+-> >> U X

=3 o o CL c cu o 1 u co

1-

maS-qj o

I— 4-

CVI

cu3

CD

to (0

M0I1YIA3Q i?J33a3d

185
2 -

1 —

t ,6 >u

Qo ,4 <

crLU

.2

.1

Longitudinal (variable temperature apparatus)

O Longitudinal

+ Transverse

(fixed point apparatus)

6 8 10 20 40 60 80 100

TEMPERATURE/ K

200 300

FIGURE 3 - Thermal conductivity of longitudinal and tra'nsverse boron-epoxy composite specimens

5*

:

186
NBS-1 14A (REV. 7-73)

U .S. DEPT. OF COMM.

BIBLIOGRAPHIC DATA

SHEET

1. PUBLIC ATION OR REPORT NO. NBSIR 75-828

2. Gov't Accession

No.

3. Recipient's Accession No.

4. TITLE AND SUBTITLE

Semi-Annual Report on Materials Research in Support of

Superconducting Machinery

5. Publication Date January 1976

6. Performing Organization ('ode 275.03

7. AUTHOR(S) R. P. Reed, J. G. Hust, M. B. Kasen, H. M. Ledbetter,

5. R. Naimon, D. T. Read, R. E. Schramm, L. L. Sparks, R. L.

8. Performing Organ. Report No. NBSIR 75-828

9. PERFORMING ORGANIZATION NAME AND ADDR ESSXobler , and W. F. Weston

NATIONAL BUREAU OF STANDARDS

DEPARTMENT OF COMMERCE

WASHINGTON, D.C. 20234

10. Project Task/Work Unit No. 2756530

11. Contract/Grant No.

12. Sponsoring Organization Name and Gomplete Address (Street, City, State, ZIP)

Advanced Research Projects Agency (ARPA)

1400 Wilson Boulevard

Arlington, Virginia 22209

13. Type of Report & Period Covered Interim

April - September 1975

14. Sponsoring Agency Code

15. SUPPLEMENTARY NOTES

16. ABSTRACT (A 200-word or less factual summary of most significant information. If document includes a significant

bibliography or literature survey, mention it here.)

Results are reported of a six-month study, March through August 1975, on candidate

materials for superconducting machinery. The results cover five areas—advanced composites,

elastic properties, fatigue resistance and fracture toughness, magnetothermal conductivity,

and thermal conductivity. Material properties were studied over the temperature range 4 to 300 K. Materials studied include: oxygen-free copper; copper-nickel alloys; a precipitation-hardening copper alloy; invar; nickel-chromium-iron alloys; stainless steels; and the

composite materials boron/aluminum, boron/epoxy, S-glass/epoxy , graphite/epoxy , and an organic-f iber/epoxy. Some notable results of the study are: the first 4 K fatigue data

on a composite material; a ten-fold increase in the fatigue life of a uniaxial glass/epoxy

composite between room temperature and liquid-hleium temperature; the first 4 K fatigue

fracture toughness studies on a nitrogen-strengthened chromiun-nickel-manganese steel,

which show this material has higher yield strength and adequate toughness compared to conventional stainless steels; room- temperature elastic properties of a copper-cadmiumchromium precipitation-hardening alloy, which are quite different from those of unalloyed

copper and show a non-parallel behavior of the shear modulus and the bulk modulus; the

thermal conductivity of 304 stainless steel may be reduced one third at 4 K by a 6 MA/m

(80 kOe) magnetic field; the first systematic study of the tensile properties of fiberreinforced composite materials between room temperature and liquid-helium temperature.

This work was supported by the Advanced Research Projects Agency of the U.S. Department

of Defense.

17. KEY WORDS (six to twelve entries; alphabetical order; capitalize only the first letter of the first key word unless a proper name; separated by semicolons

)

Composites; copper alloys; elastic properties; engineering materials; fatigue ; fracture

;

iron alloys; cryogenic temperatures; mechanical properties; nickel alloys; supercon18. AVAILABILITY |XX Unlimited

I For Official Distribution. Do Not Release to NTIS

Order From Sup. of Doc, U.S. Government Printing Office Washington. D.C. 20402. SD Cat. No. C13 . X X Order From National Technical Information Service (NTIS)

Springfield, Virginia 22151

19. SECURITY CLASS

(THIS REPORT)

UNCLASSIFIED

21. NO. OF PAGES

193

20. SECURITY CLASS

(THIS PAGE)

UNCLASSIFIED

22. Price

$7.00

USCOMM.DC 29J42-F74