IIi1ij! iiii
NBSIR 74-393

SEMI-ANNUAL REPORT ON MATERIALS RESEARCH

IN SUPPORT OF SUPERCONDUCTING MACHINERY

R. P. Reed, R. L. Durcholz, F. R. Fickett, P. J. Giarratano,

J. G. Must, M. B. Kasen, H. M. Ledbetter, R. P. Mikesell,

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

and W. F. Weston

Cryogenics Division

Institute for Basic Standards

National Bureau of Standards

Boulder, Colorado 80302

October 1974

Prepared for

Advanced Research Projects Agency

1400 Wilson Boulevard

Arlington, Virginia 22209

NBSIR 74-393

SEMI-ANNUAL REPORT ON MATERIALS RESEARCH

IN SUPPORT OF SUPERCONDUCTING MACHINERY

R. P. Reed, R. L. Durcholz, F. R. Fickett, P. J. Giarratano,

J. G. Must, M. B. Kasen, H. M. Ledbetter, R. P. Mikesell,

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

October 1974

Prepared for

Advanced Research Projects Agency

1400 Wilson Boulevard

Arlington, Virginia 22209

u s DEPARTMENT OF COMMERCE. Frederick B Dent. Secretary

NATIONAL BUREAU OF STANDARDS Richard W Roberts Director
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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, 1973 - August 9, 1974

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 of six months of study on Materials Research in Support

of Superconducting Machinery (April through September, 1974) are reported

to the sponsor, the Advanced Research Projects Agency of the U.S. Department

of Defense. The report is divided into five sections: thermal conductivity,

magnetothermal conductivity, fatigue and fracture-toughness properties, properties of advanced composites, and elastic properties. The temperature

range 4 to 300 K is covered by the study. Materials studied are either being

used or are candidates for use in superconducting machinery and include:

aluminum alloys, composites, inconels, OFHC copper, stainless steels, and

titanium alloys.

Special results of the study include: fracture-toughness and fatiguecrack growth-rate data for AISI 310, Ti-5 Al-2,5 Sn, A286 stainless steel, and

Inconel 750 at 4, 76, and 300 K; initial reports of tensile testing of composites

at 4 K; a second review paper on advanced-composite low-temperature behavior;

magnetothermal conductivity measurements on OFHC copper and Inconel 718 indicate

that, in a magnetic field, the thermal conductivity may change by 100 percent;

anomalous low-temperature elastic behavior of AISI 300 series stainless-steel

alloys; and thermal conductivity data for OFHC copper, Inconel 750, and Inco

Low-Expansion-Alloy. These data provide considerable insight into material

characteristics at extremely low temperatures, assisting in material selection

and efficient design.

iv
Report Contents

Page

Thermal Conductivity

J. G. Hust and P. J. Glarratano 1

Magnetothermal Conductivity

L. L. Sparks and F. R. Flckett 37

Fatigue and Fracture-Toughness Testing at Cryogenic Temperatures

R. L. Tobler, R. P. Mlkesell, R. L. Durcholz,

and R. P. Reed 57

Advanced Composites

M. B. Kasen and R. E. Schramm 123

Elastic Properties of Engineering Materials at Cryogenic Temperatures

H. M. Ledbetter, E. R. Nalmon, and W. F. Weston 217

Appendix A: Publication - E. R. Nalmon, W. F. Weston, and H. M.

Ledbetter, Elastic properties of two titanium alloys at low

temperatures, Cryogenics 14 (1974) 246-9 27 3

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 re- ported 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

THERMAL CONDUCTIVITY

J. G. Hust and P. J. Giarratano

Cryogenics Division

Institute for Basic Standards

National Bureau of Standards

Boulder, Colorado 80302

October 1974
Summary: Thermal Conductivity

Thermal conductivity, electrical resistivity, thermopower, and

Lorenz ratio data have been obtained for OFHC copper, Inconel X-750, and

INCO LEA (Low Expansion Alloy) . Values of these properties are presented in graphical and tabular form and equations are provided for con- venient computer calculations for data at temperatures from 4 to 300 K.

The thermal conductivity data are accurate to 2%, which is well within

material variability limits expected for production heats. Prediction

techniques have been shown to yield thermal conductivity values within

15 percent of experimental values for pure metals, such as OFHC copper.

In addition, thermal conductivity values have been obtained, using

the fixed-point apparatus, for Kromarc-58 and AISI-310 stainless steels

at 6, 79, 196, and 277 K. These values, accurate to 10%, are presented in

graphical and tabular form.

Contents; Thermal Conductivity

Page

1. Introduction 3

2. Apparatus 4

3. Specimen Characterization 5

4. Results 6

5. Discussion 7

6. Acknowledgements 9

7. References 10

List of Figures 11

List of Tables 30

2
THERMAL CONDUCTIVITY

J. G. Hust and P. J. Giarratano

1, Introduction

Optimum design of superconducting machinery requires a knowledge

of thermal properties of technically important materials at low temperatures. But, often, components extend from ambient temperatures to low

temperatures, and, therefore, data are needed from 4 K to 300 K. Thermal

conductivity of metals is a function of the host metal and its constituent

elements and temperature; but it also depends on the microstructure of

the material, strain, fatigue, and heat treatment; and in some cases environ

mental effects are important (e.g., in the case of superconducting machinery

magnetic field intensities may produce significant changes). Specific

thermal conductivity data for particular components are required to limit

heat losses and thermal stresses in machines. With such a myriad of

parameters it is difficult (at best) to determine accurately, by experiment, the thermal conductivity of each machine component. Instead, we

often rely on a limited number of measurements on specific types of

materials, characterized according to the above parameters. Due to

material variability effects these values are generally considered

accurate to only about 10%.

It should be noted that accurate thermal conductivity measurements

are much more expensive and time consuming relative to many other property measurements. The ease with which high accuracy can be obtained

is primarily a function of temperature and the conductivity itself. The

extremes of conductivity (both very high, such as in pure metals, and

very low, such as in insulating materials) promote inaccuracies in

measurement. Generally, low temperature data (below about 100 K) are

more readily obtained with accuracy because radiation errors can be

minimized

For the above reasons, various methods of obtaining thermal con- ductivity data are used. The method used depends primarily on the

accuracy dictated by the application. Usually accuracies of 10% are sufficient; in other cases values accurate to 50% suffice. The present state-of-the-art in thermal conductivity measurements is about

3
1% accuracy, but seldom is this required except in theoretical or standardization work.

In this laboratory three methods of data determination are used:

(a) A variable-temperature apparatus, capable of measuring at any

temperature from 4 to 300 K, with an accuracy of about 2%. This

apparatus also simultaneously measures electrical resistivity and

thermopower for additional specimen characterization and subsequent

predictions of similar materials. This apparatus will be used to

measure materials that have not been sufficiently measured and

characterized previously and for accurate temperature dependence

determinations.

(b) A fixed point apparatus, capable of measuring only near the fixed

temperatures of various boiling fluids (such as liquid helium and

liquid nitrogen) and melting or subliming solids (such as ice and

CO2) . The accuracy of this apparatus is about 10%, depending

mainly on the conductivity of the specimen.

(c) Predictive methods such as those described by Hust and Clark

[1] and Hust and Sparks [2]. The latter paper also contains data

for many metals and alloys, necessary to carry out such predictions.

2 . Apparatus

The variable-temperature apparatus was used for the measurements

reported for Inconel X-750, OFHC copper, and INCO LEA (Low Expansion

Alloy). This is a multi-property apparatus used to simultaneously measure thermal conductivity, electrical resistivity, and thermopower at temperatures between 4 and 300 K. This apparatus is based on the longitudinal

heat flow method. The specimen is a rod 23 cm long with a diameter up to

3 cm depending on its conductivity. The top of the specimen is thermally

connected to a temperature-controlled floating sink and the gradient heater

is connected at the bottom of the specimen. The specimen is surrounded with

insulation and a temperature-controlled concentric shield. The longitudinal

temperature distribution of the shield is adjusted to match that of the specimen.

Eight equally spaced thermocouples are attached to the specimen to

determine its longitudinal temperature distribution. These thermocouples are calibrated in-place using accurately calibrated platinum and

germanium resistance thermometers. Further details of this apparatus
are documented in an unpublished report available from Hust. A schematic diagram of this cryostat is shown in the previous semi-annxial

report to ARPA [3].

The fixed-point apparatus was used to measure Kromarc-58 and

AISI-310 stainless steels. This longitudinal-heat-flow apparatus is

designed to perform measurements of 10% accuracy with greater speed and

smaller specimen size than the multiproperty apparatus. The specimen is a 5 cm long rod with a diameter up to 3 cm depending on its conductivity.

Further details of this apparatus are available in an unpublished report

from Hust. A schematic diagram of this cryostat is shown in the previous

semi-annual report to ARPA [3]

.

3. Specimen Characterization

The materials investigated have been selected from a long list of

materials that have applications in superconducting machinery. The

materials considered here are OFHC copper, Inconel X-750, INCO LEA (Low

Expansion Alloy), Kromarc-58 stainless steel, and AISI-310 stainless steel.

OFHC copper was chosen as the first specimen to be measured, not only because of its importance in superconducting machines, but also because a good conductor was needed to perform initial checks and thermocouple calibrations of the apparatus.

The OFHC copper specimen was fabricated from production copper as follows: A 6.4 mm copper rod was degreased with freon and etched with

a solution of 50% water and 50% nitric acid. The rod was swaged to

3.26 mm with cleaning between each swaging step and vacuum anneals

whenever flaking became evident. The final anneal before measurement

was 650°C for one hour. Prior to each aimeal the specimen was acid

etched.

The Inconel X-750, INCO LEA, Kromarc-58 stainless steel, and AISI-310

stainless steel stock material was obtained from Westinghouse Electric Corporation, Research and Development Center and is the same material being tested

by them in this program. The Inconel X-750 was in the solution treated (ST)

condition (1800^ - one hr - air cooled) . The INCO LEA was in the solution

treated aged (STA) condition (1800°F - one hr - air cooled - 1225 "F - 8 hrs - air cooled) . Both of these materials were machined to cylindrical rods 23 cm

long and 11.3 mm diameter with no subsequent heat treatment.

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The Kromarc-58 stainless steel was annealed at 1800 °F for one hour, ; water quenched, and then cold worked 30% in the transverse direction. The

| AISI-310 stainless steel was heated to 2000°F for one hour and water quenched.

Both of these materials were machined to cylindrical rods 5 cm long and 11.3

ram diameter with no subsequent heat treatment.

4. Results

Measurements were performed on specimens of OFHC copper, Inconel

| X-750, and INCO LEA (Low Expansion Alloy) over the temperature range

4 to 300 K. Fifty-eight runs were conducted on OFHC copper, 29 on Inconel X-750, and 28 on INCO LEA. OFHC copper was studied more ex- tensively than the other two materials to establish system characteristics

and to perform an extensive thermocouple calibration data set. The following equations were least-squares fitted to these data using orthonormal ! fitting techniques:

enX

P

S

where k = thermal conductivity, p = electrical resistivity, S = thermopower, and T = temperature. Temperatures are based on the IPTS-68

scale above 20 K and the NBS P2-20 (1965) scale below 20 K. The parameters, a^, h^t and c_j^ determined by least squares, are presented in

table 1. Further details of this procedure are described by Hust, et al.

[4]. The deviations of the experimental data from these equations

are given in figures 1 through 9. The horizontal bars in figures 2, 3,

5, 6, 8 and 9 indicate the temperature span across the specimen for each

run. The "observed" thermal conductivities are computed from the mean temperature gradients indicated by adjacent thermocouples. Calculated

values of X, p, S, and L = pX/T (Lorenz ratio) are presented in tables

2, 3, and 4 and figures 10 through 21.

nS aJCnT]

T i+1

i+1

(1)

m

I b.[CnT]

i=l ^ i-1

(2)

S c.[CnT' l^'/r ; T = + 1

1=1

(3)

6
A detailed error analysis for this system has been performed and

is available as an unpublished report from Hust. Based on this analysis

of systematic and random errors, typical uncertainty estimates (with

95% confidence) are as follows:

thermal conductivity: 2% from 4 to 20 K, 1% from

20 K to 200 K, and 2% from

200 K to 300 K

electrical resistivity: 0.2%

thermopower: 0.1 [iV/K up to 20 K

and 0.01 ^V/K above 20 K

The actual uncertainties depend on the conductance of the specimen.

For Inconel X-750 and INCO LEA the t5T)ical values are applicable. Since

the conductance of the copper specimen was significantly larger, especially

at liquid-helium temperature, the uncertainty for copper is larger at low

temperatures. The uncertainties in thermal conductivity and electrical re- sistivity near 4 K are 4% and 1%, respectively, decreasing to the typical

values above 20 K.

The thermopower values given here are absolute values although

our measurements were carried out with respect to normal silver wire.

The absolute thermopowers of normal silver reported by Borelius, et

al. [5] were used to convert the experimental data to the absolute scale.

Thermal conductivity measurements were performed on Kromarc-58 and

AISI-310 stainless steels at 6, 79, 196, and 277 K. The results of these

measurements are listed in Table 5 and illustrated in figure 22. The smooth

curves shown in figure 22 are based on the temperature dependence of similar

alloys. The measured values of thermal conductivity have been compared to

the fee steel data tabulated by Childs et al. [6] and are in good agreement

with the band of values compiled.

5. Discussion

OFHC copper is a relatively pure conmercial copper and, therefore,

its low- temperature thermal conductivity and electrical resistivity are highly variable from specimen-to-specimen. This variability is due to

the fact that the low- temperature conductivity is nearly totally determined by the imperfections (impurities and physical defects) in the

7
metal. Small absolute variations in impurities represent large relative

variations in pure metals. Fortunately, it is relatively simple to

characterize the degree of imperfections in a pure metal by electrical

resistivity data, which is much easier to obtain than thermal conductivity data. Often the residual resistivity or the residual resistivity

ratio is used to indicate the effective electrical purity of the specimen.

For all but the most pure research materials the resistivity at 4 K

is the residual resistivity. The residual resistivity ratio is the

ratio of resistivity at near ambient conditions to the residual resistivity.

The resistivity at ambient temperature is a characteristic of the host metal

and the residual resistivity is characteristic of its imperfections, both

chemical and physical. We normally use the resistivity at the ice-point for

defining this ratio, RRR = P273 15k''P4K* Other OFHC copper specimens that

we have characterized at this laboratory have RRR values ranging from about

50 to 200. The particular specimen reported on here has an RRR of

250, somewhat more electrically pure than normal. This may be due to

lower chemical impurity level or to impurity oxidation during the annealing

process.

As indicated in the previous semi-annual report on this project [3],

thermal conductivity of a pure metal can be predicted reasonably accurately

from a knowledge of the Lorenz ratio, residual resistivity, and Matthiessen's

rule [1,2]. Based on this information values of thermal conductivity

for OFHC copper were presented in the last report. The current measurements

show that those predictions were accurate to 10% above 15 K. At 10 K

the difference between measurement and prediction is 15%. This exercise

serves to illustrate the power of predictive methods for pure metals

whose low- temperature conductivities may range over several orders of

magnitude. At the lowest temperatures where this method was least accurate,

other methods can be applied to yield more accurate predictions, but

these are not discussed here.

The thermal conductivity of highly alloyed structural metals is not

readily predicted. This is primarily because in such alloys the thermal

transport is by quantized lattice vibrations (phonons) as well as by

the electrons. The phonon conductivity is often much larger than the

8
electronic conduction and, thus, the electrical resistivity is no longer

directly linked with thermal conductivity through the Lorenz ratio.

Inconel X-750 and INCO LEA are part of a class of materials that have

very high phonon-to-electronic conductivity at lower temperatures and

therefore, their conductivities cannot be predicted from the Lorenz

ratio. Lorenz ratios above the classical Sommerfeld value (2.443 x

_g 2 2

10 V /K ) indicate the presence of significant phonon conduction.

The low- temperature Lorenz ratio of Inconel X-750 indicates a very high

relative phonon component (about ten times greater than the electronic

component) while for INCO LEA the phonon component is about twice as large as the electronic component.

6. Acknowledgements

We wish to express our appreciation to Joe Wells of Westinghouse

Research Laboratories, Pittsburgh, Pa. for the material which he so generously provided for these measurements.

9
7. References

1. Hust, J. G. and Clark, A. F., The Lorenz ratio as a tool for predicting

the thermal conductivity of metals and alloys. Materials Research

and Standards 11, (No. 8), 22-4 (Aug. 1971).

2. Hust, J. G. and Sparks, L. L., Lorenz ratios of technically important

metals and alloys, Nat. Bur. Stand. (U.S.) Tech. Note 634, (Feb.

1973) 133 pp.

3. Sparks, L. L., Fickett, F. R. , Hust, J. G., et al.. Semi-annual

report on materials research in support of superconducting machinery,

NBSIR 74-359 (1974) 308 pp; available from NTIS, AD 780-596/3WM.

4. Hust, J. G., Weitzel, D. H., and Powell, R. L., Thermal conductivity

standard reference materials from 4 to 300 K, I. Armco iron: Including apparatus description and error analysis, J. Res. NBS

74A, 673-90 (1970).

5. Borelius, G., Keesom, W. H. , Johansson, C. H., and Linde, J. 0.,

Establishment of an absolute scale for the thermoelectric force,

Proc. Kon. Akad. Amsterdam 35, 10 (1932).

6. Childs, G. E., Ericks, L. J., and Powell, R. L., Thermal conductivity

of solids at room temperature and below: A review and compilation of

the literature, Nat. Bur. Stand. (U.S.) Monograph 131 (Sept. 1973)

608 pp.

10
List of Figures

Page

1. Thermal conductivity deviations for OFHC copper 12

2. Electrical resistivity deviations for OFHC copper . 12

3. Thermovoltage deviations for OFHC copper 13

4. Thermal conductivity deviations for Inconel X-750 13

5. Electrical resistivity deviations for Inconel X-750 14

6. Thermovoltage deviations for Inconel X-750 14

7. Thermal conductivity deviations for INCO LEA 15

8. Electrical resistivity deviations for INCO LEA „ . « . , 15

9. Thermovoltage deviations for INCO LEA » . 16

10. Thermal conductivity of OFHC copper 17

11. Electrical resistivity of OFHC copper 18

12. Thermopower of OFHC copper . 19

13. Lorenz ratio of OFHC copper 20

14. Thermal conductivity of Inconel X-750 21

.15. Electrical resistivity of Inconel X-750 22

16. Thermopower of Inconel X-750 23

17. Lorenz ratio of Inconel X-750 24

18. Thermal conductivity of INCO LEA 25

19. Electrical resistivity of INCO LEA 26

20. Thermopower of INCO LEA 27

21. Lorenz ratio of INCO LEA 28

22. Thermal conductivity of Kromarc-58 and AISI-310 stainless steels . . 29

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8 +

I 1 I I I 60 ISO 2iO 323 TEMPERATURE (Kelvin)

Figure 1. Thermal conductivity deviations for OFHC copper.

TEMPERATURE (KtU In)

Figure 2. Electrical resistivity deviations for OFHC copper.

12
60 160 240 320

TEMPERATURE (KtU In)

Figure 3. Thermovoltage deviations for OFHC copper.

1 <•

z

2 0 >

UJ . o -I

UJ

a: -2 ui

a.

-3"

/ + + +

\ * ^ t

* * ^

.

*

*

1 L

+ +

60 160

I I 240 320

TEMPERATURE (Ktlv I n;

Figure 4. Thermal conductivity deviations for Inconel X-750.

13
z

o

>

Hi

o

UJ

Q.

Figure 5,

eO 1(0 240 TEMPERATURE (KtUU)

Electrical resistivity deviations for Inconel X-750.

ID ICO 240

TEMPERATURE (KtltfU)

Figure 6. Thermovoltage deviations for Inconel X-750.

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1

I 1 1 +

+

> 1 > 1

• *

-— —.

^* -

t'^ .

^ ^

T * *

1 1 1 » eo 160 240 320 TEMPERATURE (K*lv I h)

Figure 7. Thermal conductivity deviations for INCO LEA.

eS ISO 2iO 320 TEMPERATURE (KeU I hi

Figure 8. Electrical resistivity deviations for INCO LEA.

15
c

Figure 9.

TEMPERATURE (KeUU)

Thermovoltage deviations for INCO LEA.

16
60.1 I 1 I t

I II > I > I I

I I

I >

6 e :G 20 40 60 80 100 200 40 TEMPERATURE (KtU In)

Figure 10. Thermal conductivity of OFHC copper.

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TEMPERATURE (K« Win)

Figure 11. Electrical resistivity of OFHC copper.

is
2.3 I

I I I t I i

80 160 240 320 TEMPERATURE (Ktlv In)

Figure 12. Thermopower of OFHC copper.

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2.6 r

TEMPERATURE (Ktlv I h)

Figure 13. Lorenz ratio of OFHC copper.

20
20 I

I < I » » tI » I t I t » I I t I > I

' I ' i ' < ' I ' <

I i

I I I I I i 'I,

' i 6 e 10 20 iO GO 80 100 200 4O0 TEMPERATURE (KtUlh)

Figure 14. Thermal conductivity of Inconel X-750.

21
1.29 I

» I I I I I

TEMPERATURE (K« I vlh)

Figure 15. Electrical resistivity of Inconel X-750.

22
23
24
25
26
TEMPERATURE (Kelvin)

Figure 20. Thermopower of INCO LEA.

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28

List of Tables

Page

1. Parameters in equations 1, 2, and 3 31

2. Thermal and Electrical Properties of OFHC copper 32

3. Thermal and Electrical Properties of Inconel X-750. ...... 33

4. Thermal and Electrical Properties of INCO LEA ......... 34

5. Thermal Conductivity Data for Kromarc-58 and AISI-310

Stainless Steels 35

30
Parameters in equations 1, 2, and 3.

OFHC COPPER (I)

THERMOPOWER

-3.9fl7e4822t002 I.84739120»003

"3.5383707lffl03

3. 39074429t003

-1.7G743497t003

5. 10924193*002

-7.7l5l8S43t001

4. 7615lS07t000

INCONEL X-750 SKMPll

ELECTRICAL

CONDUCTIVITY RESISTIVITY THERHOPO^ER

I -G.45I63672t00I 1.21912799 -OOG -2.2532403Gt001

2 I.82907S93t002 4.51454338 -008 1.4302G970t002

3 -2. 32001559*002 -3. 008G0094 -008 -1.85G035I4t002

4 I.726303G3f002 8. G4270587 -009 1.07305G20f002

5 -8.3l7G929flt001 -I. 19334919 -009 -2. 49926242*001

G 2. 70eOGGI2tOOI 7. 1525G039 on 2.93902707tOOO

7 -G. 03I9G69etOOO

e 9. 07810755-001

9 -8. 83870489-002

10 5. 030II99G-003

II -I.2714G5G9-004

CONDUCTIVITY RESISTIVITY THERHOPOITER

-:. :7912I20tOOI 9. 97525213 -007 3.30951242«002

2 :. 93895753*001 -7. 51 11 1280 -007 -I.85240704t003

3 -:.405«059et00I 7. 81142520 -007 3.95194962*003

I 5. 73e03284tOOO -4. 39615174 -007 -4.43311928*003

5 -:. 40547958*000 1. 45204569 -007 2.74953751*003

G 2. 051438i7-OOI -2. 81979266 -008 -9.GI465042t002

-:.G473237I-002 2. 98050307 -009 1.75464330*002

8 5.G099I387-004 -I. 30724755 -010 -1.26795595*001

9 -:.07G58075-003 -1. 85507554 -009

mm 5. C3Dn99«-D03 1. 20856094 -010

-:. 27:46569-004 -3. 45614932 -012

INCO L£* STA

THEW1AL ELECTRICAL

31
Table 2. Thermal and Electrical Properties of OFHC copper.

TEMP THERMAL ELECTRICAL CONDUCTIVITY RESISTIVITY NUHKR

(K) [\i/yi Ki mno om hi (V/K)2X10t-8) IMICROVOLTS/Kl

e 2730 0.07156 2.45 •0.21 9 3000 0.07302 2.44 •0.17

10 3230 0.07305 2.36 •0.16

12 3610 0.07132 2.15 -0.16

14 3900 0.07070 1.97 •0.16

16 4070 0.07243 1.84 •0.14

le 4120 0.07605 1.74 •0.10

20 4050 0.08098 1.64 •0.03

25 3550 0.09911 1.41 0.19

10 2870 0.1324 1.27 0.46

35 2250 0.1915 1.23 0.71 40 1770 0.2845 1.26 0.91 45 1410 0.4146 ;' 1.30 1.06

50 1150 0.5819 f 1.34 1.15

60 834 1.017 1.41 1.25

70 664 1.561 1.46 1.26

80 568 2.160 1.55 1.23

90 513 2.849 1.62 1.19

100 480 3.545 1.70 1.15

120 448 4.972 1.86 1. 10 140 434 6.397 1.98 1.11 160 426 7.801 2.08 1.15

180 420 9.185 2.14 1.21 200 415 10.56 2.19 1.30 220 412 11.92 2.23 1.39

240 411 13.28 2.28 1.50 260 412 14.64 2. 32 1.60 280 412 15.99 2.35 1.71 300 411 17.32 2.37 1.82

32
Table 3. Thermal and Electrical Properties of Inconel X-750.

TEW> THCTHAL ELECTRICAL LORENZ THER«OPO«R

C0»OUCTIVITY RESISTIVITY NUHBER

(K) («(/H K) [NANO OHH HI (V/K}2X10(-ei (HICROVOLTS/Kl

6 0.962 1242 19.92 0.41 7 1.18 1242 20.92 0.49

8 1.39 1241 21.55 0.55

9 1.60 1241 22. 05 0.60

10 1.81 1240 22.51 0.65

12 2.26 1239 23.29 0.76

U 2.70 1239 23.87 0.85

IG 3.13 1238 24.21 0.94

18 3.54 1238 24.35 1.02

20 3.93 1238 24.34 1.08

25 4.82 1238 23.84 1.23

30 5.58 1238 23.01 1.38

35 6.23 1238 22. 04 1.54

iO 6.79 1238 21.03 1.70 45 7.27 1238 20.00 1.87

50 7.66 1239 18.99 2.03

60 8.26 1240 17. 07 2.35

70 8.66 1241 15.35 2.67

80 8.93 1242 13.87 2.98

90 9.12 1243 12.60 3.27

100 9.27 1244 11.54 3.55

120 9.50 1247 9. 88 4.09

liO 9.71 1250 8.67 4.58

160 9.92 1252 7.77 5.03

180 10.1 1255 7.07 5.45

200 10.4 1257 6.53 5.84

220 10.6 1260 6. 09 6.20

2iO 10.9 1263 5.75 6.54

260 11.2 1265 5.46 6.86

280 11.5 1266 5.22 7.16

30C 11.8 1270 5.00 7.45

33
Table 4. Thermal and Electrical Properties of INCO LEA.

TEHP THERHAL ELECTRIC/U. LOREKZ THEWOPOiER

CONDUCTIVITY RESISTIVITY NUMBER

IKl IM/n K) IV/K]2X10[-e] ihicrovolts/k:

6 0.297 697.4 3.46 -0.37

7 0.361 697.6 3.60 -0.62

8 0.426 698.1 3.72 -0.80

9 0.490 698.4 3.80 -0.95

:fl 0.554 698.6 3.87 -1.06

12 0.683 699.1 3.98 -1.22

li 0.814 699.5 4. 07 -1.37

IC 0.948 700.0 4. 15 -1.55

18 1.09 700.6 4.23 -1.75

20 1.23 701.2 4.32 -1.97

25 1.61 703.0 4.52 -2.56

30 2.00 705.2 4. 69 -3.11

35 2.39 707.7 4.83 -3.60

40 2.78 710.4 4.94 -4.03

45 3.16 713.5 5.00 -4.41

50 3.52 716.8 5.05 -4.76

SO 4.20 724.1 5. 06 -5.36

73 4.81 732.1 5.03 -5.87

80 5.37 740.9 4.97 -6.30

90 5.88 750.2 4. 90 -6.66

:oo 6.36 759.9 4.83 -6.95

120 7.22 780.3 4. 69 -7.32

:4o 7.97 801.4 4.57 -7.45

:co 8.65 822.9 4.45 -7.37

180 9.24 844.3 4.34 -7.13

200 9.77 865.5 4.23 -6.76

220 10.2 886.1 4.13 -6.31 240 10.7 906.1 4. 04 -5.80 2C0

« « • 925.4 3.96 -5.26

280 ii'6 943.9 3.90 -4.72

3s: 12.

1

961.4 3.87 -4.17

34
Table 5. Thermal conductivity data for Kromarc-58 and

AISI-310 stainless steels

Material

Kromarc-58

AISI-310

Temperature (K)

6. 50

79. 1 196

277

6.46

79.4

196

277

X(W/mK)

0. 440

5. 90

9. 70

11. 5

0. 720

7. 15

10. 8

12. 5

35
36
NBSIR

SEMI-ANNUAL REPORT ON MATERIALS RESEARCH IN SUPPORT OF

SUPERCONDUCTING MACHINERY

MAGNETOTHERMAL CONDUCTIVITY

L. L. Sparks and F. R. Fickett

Cryogenics Division

Institute for Basic Standards

National Bureau of Standards

Boulder, Colorado 80302

October 1974
Summary; Magnetothermal Conductivity

The magneto thermal conductivities of Inconel 718 and oxygen-free

copper have been determined. The data are presented as thermal conductivity as a function of temperature, A-(T), thermal conductivity as a function

of magnetic field, X(H) , and as the relative change in thermal resistivity as

a function of magnetic field, AW/W^ (H) . When possible the data are compared

to existing data. In these limited instances, agreement is within the accuracy

of our measurements. The effect of a magnetic field is always to reduce the

thermal conductivity. The reduction, at 8 Tesla, for Inconel 718 is approximately 15% at 5.5 K and 50% at 19.5 K; for copper (RRR = 107), the reduction

is about 100% at 5.5 K and 60% at 19.5 K. The magnetic-field effect increases

with increasing temperature for Inconel and decreases with increasing temperature for copper. The temperature dependence of AW/W^ exhibited by Inconel

is associated with the negative temperature derivative of the zero field re- sistivity, dp /dT, shown by Inconel 718 below 20 K.

Contents; Magneto thermal Conductivity

Page

1. Introduction 39

2. Procedures 39

2.1 Apparatus 39

2.2 Methods 40

2.3 Materials 41

3. Results 41

4. Discussion 42

5. References 45

List of Figures 46

38
MAGNETOTHERMAL CONDUCTIVITY

L. L. Sparks and F. R, Picket

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, et al. [1] in the previous report.

The materials studied in this program are being used in superconducting

motors and generators. Optimum design of these machines, which must operate

at temperatures between 4 K and 20 K while in magnetic fields up to 10 Tesla,

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 Inconels, stainless steels, and

metallic composites.

The experimental determination of X(H) of metals in high fields is com- plicated 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 gradients along the

specimens. Our X(H) probe is designed to be used in a superconducting solenoid

with a 3.8 cm bore and a 2.5 cm homogeneous (1%) field sphere. Our specimen

length is therefore limited to approximately 2.5 cm.

2. Procedures

2 . 1 Apparatus

A brief review of the cryostat, Instrumentation, and experimental procedures is given below. For further details refer to [1].

Figure 1 shows the principal components of the probe with a specimen

mounted in the longitudinal position (magnetic field parallel to the heat

flow) . A shorter specimen may be mounted horizontally so that the heat flow

is perpendicular to the magnetic field for transverse measurements.

39
The cooling required to attain specimen temperatures in the range

between 4 and 20 K is provided by immersing the probe chamber in liquid

helium. Radiative heat transfer is minimal in this temperature range, and

gas conduction is eliminated by evacuating the specimen chamber to 5 x 10

^

torr. Our operating pressure is reduced to less than this by a gas adsorption agent in the sealed probe space.

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 are 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 eliminate any temperature difference between

the specimen and the tempering post. 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, which are clamped to the specimen. The GRT is located in the TCHS

and is used for zero-magnetic-field calibration of the GRT resistors. The

effect of the magnetic field on the resistance of the CRT's is taken to be

that published by Neuringer and Shapira [2]. All 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.

2.2 Methods

Determination of X(H) and the electrical magnetoresistivity, p (H)

,

requires two types of tests, which are referred to as isothermal tests and

gradient tests. For an isothermal test the CRT resistances are read with

the specimen heater off. A series of these measurements at different temperatures determines the resistance-versus-temperature characteristics for

40
each CRT at zero field. Data from Neurlnger and Shapira [2], on values

of (R^ - Rq)/Rq then allows the CRT's to be calibrated for H # 0. Gradient

tests are made with the SPECIMEN HEATER on. This power establishes a thermal gradient along the specimen. Measurement of the heater power and

the resulting CRT resistances allows the specimen thermal conductivity to

be determined. Electrical resistance is determined by passing a current

through the specimen and measuring the resultant voltage drop. The resistivity

is calculated at the average temperature of the specimen.

2.3 Materials

Two materials have been tested - Inconel 718 and oxygen-free copper.

The Inconel specimen was in the age-hardened condition. The Rockwell hardness

was C39 (not B39 as reported in [1]). The copper specimen was cold swaged

from 0.476 cm diameter rod to 0.25 cm diameter rod. The rod was etched and

annealed several times during the swaging process. The final anneal was at

—6

850°C for one hour in a vacuum of 10 torr. The resistivity of the resulting

specimen at 5 K is 0.146 nAn (RRR = 107).

3. Results

The first material to be tested in the magnetothermal conductivity

system was Inconel 718. A specimen from the same heat was tested by Hust,

et al. [3]; zero field values of X and p were determined. Their experiment,

done in the NBS precision thermal-conductivity apparatus, has an estimated

error of not more than 2%. Their results for A,(T) are labeled HUST on figure

2. The figure also shows our data on thermal conductivity as a function of

temperature at several magnetic fields. Comparison of electrical resistivities

between the two specimens at zero field indicates a difference of about 0.5%.

Figures 3 and 4 present thermal conductivity and relative change in thermal

resistance, AW/W^, as functions of magnetic field with temperature as a parameter.

The ordinate in figure 4 is equivalent to (A, - X ) /X using the general relao H H

tionship W(thermal resistivity) = Figure 5 gives thermal conductivity as

a function of temperature with magnetic field as a parameter for our copper

specimen. Figures 6 and 7 present thermal conductivity and relative change

in thermal resistance as functions of magnetic field with temperature as a parameter

.

41
At present the best estimate of the system errors Indicates probable

uncertainty of 10 - 15% in the thermal conductivity. Comparison of our zero-field data to that of Hust, et al. [3] for Inconel 718 indicates

Z^A(average) = 8.6% at 7 K, 5.7% at 12 K, and 4.6% at 19 K. These figures

include possible real differences between the two specimens. One possible

source of error that will be eliminated in the future is the value of the

magnetic field in our solenoid. At the present time the magnetic-field —

magnet- current relationship is uncertain by approximately 2 - 6%. We re- cently acquired a precision rotating-coil gaussmeter that will allow calibration of the current-field relationship to approximately 0.1%. If, in

fact, our present field calibration is significantly changed, both the

Inconel 718 and the copper data will be appropriately adjusted.

4. Discussion

Thermal conductivity is a complex phenomenon because many separate

mechanisms are operating simultaneously to transmit energy through the

material. Thermal conductivity is defined by the heat transfer equation

g-Mf (1)

where Q is the heat current in watts (W) , A is the cross sectional area in

2 m , T is in Kelvin (K) , and thus A, has the units of W/mK.

Powell [4] gave an overview of transport properties in metals and a useful list of references. Briefly, the thermal energy in most materials is trans

ported both by the lattice vibrations (phonons) and by the conduction electrons

It is conventional to designate the lattice component A, , and the electronic

component A,^. These conductivities make up the total thermal conductivity,

k = k + k . (2) e g

Thermal resistance to each of these modes of heat flow arises from several

interactions. For electrons the principal interactions retarding heat flow

are electron-phonon, electron-impurity , and electron-defect; for phonons,

the interactions are phonon-electron , phonon—point-defect, and phonon-dislocation. The lattice contribution, A. , at low temperatures is generally less

42
than 20% of the total thermal conductivity in pure metals but may predominate for some alloys [5].

One of the classical methods used to separate and is to apply a magnetic field. This has the effect of decreasing A,^ while presumably not

affecting A, [6], One would expect that X. (total) of a pure metal would be

affected more than X (total) for an alloy. This is borne out by our data as discussed below, and by the results of Fevrier and Morize [7]; this work will

be discussed in more detail below.

The effect of an 80 kOe magnetic field is to decrease X(H) for Inconel

718 by 14% and for copper by 104% at 5.5 K. At 19.5 K the reduction is 49%

for Inconel and 58% for copper. The temperature dependence of AW/W^ for

Inconel, figure 4, is opposite that for copper, figure 7. The slope of

AW/W^(T) is positive for Inconel and negative for copper. This "anomaly"

appears to be a natural consequence of the thermal equivalent of Kohler's

rule rather than a manifestation of the ferromagnetism seen in Inconel at

these temperatures. Kohler's rule is the classical equation of electrical

magnetoresistance and can be written

^o ^o '^o

i.e., the temperature dependence of ^/p^ arises only through p^. The ferromagnetic behavior of Inconel 718 below 77 K is a factor in the temperature

dependence of AW/W , but only to the extent that it affects W . o o Figure 8c indicates the different temperature dependences discussed

above. Non-magnetic metals with no magnetic impurities show a low temperatureresistivity versus temperature curve like that shown for copper in figure 8a;

magnetic materials, on the other hand, frequently show a negative dp^/dT below

some critical temperature, as is the case with Inconel [3] (figure 8a). Since

^ (= p no- p ) is not a strong function of temperature, Ap/p (T) increases

O

when p becomes smaller (Inconel), and decreases when p becomes larger (copper)

o o as shown in figure 8b. There is no reason to expect vast differences between

the electrical and thermal resistivities, although they are not identical

functions of temperature. The thermal data illustrated in figure 8c seems

quite reasonable for both metals.

43
The relationship between the electronic contribution to the thermal

conductivity and the electrical conductivity is expressed by the WiedemannFranz-Lorenz law

(4)

Procedures for predicting X at H = 0 from L and p data are reasonably

successful and are discussed by Hust, et al. [5]. It is hoped that progress

can be made in predictive schemes based on the electrical and thermal con- ductivities in non-zero fields, i.e..

Magnetothermal data on many alloy groups will be necessary before any such

schemes can be evaluated. For the present, let it suffice to observe, in

figure 9, that the effect of increasing H is to reduce L.

In the case of copper, meaningful comparison to other data is difficult

even at zero magnetic field. The electrical and thermal conductivity of pure

copper is very sensitive to strain, thermal history, and trace impurities.

Fickett [8] discussed the large effect annealing in a partial pressure of

oxygen has on the RRR of copper. Fevrier and Morize [7] have reported relative magnetothermal conductivity data for two copper specimens, RRR = 62 and

RRR = 162. Figure 10 shows the ratio of A,(H)/X(0) as a function of magnetic

field for their two specimens and our specimen (RRR = 107) . The temperature

of the Fevrier and Morize data is not given in reference [7] but is assumed

to be 4 K; the temperature of our measurement is 5.5 K. The general behavior

of two sets of data are similar. Resistivity, and therefore Lorenz-ratio

data, are not available for our copper specimen due to the low- level voltages

involved.

L(H) = P(H)X(H)

(5) T

44
5. References

1. 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. Neuringer, L. J. and Shapira, Y., Low temperature thermometry in

high magnetic fields, I. Carbon resistors RSI 4£, (No. 10), 1314

(October, 1969).

3. Hust, J. G. , Weitzel, D. H. and Powell, R. L. , Thermal conductivity,

electrical resistivity, and thermopower of aerospace alloys from

4 to 300 K, J. Res. NBS (U.S.) (Phys. and Chem. ) 75A (No. 4), 269-277

(July - August, 1971).

4. Powell, R. L. , Thermophysical Properties of Metals at Cryogenic

Temperatures , ASTM Special Publication No. 387 , 134 pp (ASTM, New York,

1966).

5. Hust, J. G. and Sparks, L. L. , Lorenz ratios of technically important

metals and alloys, Nat. Bur. Stand. (U.S.) Tech. Note 634 (February,

1973) 127 pp.

6. van Witzenburg, W. and Laubitz, M. J. , Magnetoresistances and the phonon.

Can. J. Phys. 46, 1887 (1968).

7. Fevrier, A. and Morize, D., The effect of magnetic field on the thermal

conductivity and electrical resistivity of different materials. Cryogenics

13, 603 (1973).

8. Fickett, F. R. , Oxygen annealing of copper: a review. Mater. Sci. Eng.

14, 199 (1974).

45
List of Figures

Page

1. Magnetothermal conductivity probe and magnet 47

2. Thermal conductivity of Inconel 718 as a function of

temperature at several magnetic fields 48

3. Thermal conductivity of Inconel 718 as a function of

magnetic field at several temperatures 49

4. Relative change in thermal resistance of Inconel 718 as a function of magnetic field at several temperatures 50

5. Thermal conductivity of copper as a function of temperature

at several magnetic fields 51

6. Thermal conductivity of copper as a function of magnetic

field at several temperatures 52

7. Relative change in thermal resistance of copper as a function

of magnetic field at several temperatures 53

8. Characteristic curves of p (zero field electrical resistivity),

Ap/p (relative electrical resistivity change in a magnetic field),

and AW/W (relat ive thermal resistivity change in a magnetic

field) for Inconel and copper below 20 K 54

9. Lorenz ratios of Inconel 718 as a function of temperature at

several magnet fields 55

10. Comparison of the present magnetothermal conductivity data

to that of Fevrier and Morize 56

46
Figure 1. Magnetothermal conductivity probe and magnet.
48
Figure 3. Thermal conductivity of Inconel 718 as a function of magnetic

field at several temperatures.

49
100

H (kOe)

Figure 4. Relative change in thermal resistance of Inconel 718 as a function

of magnetic field at several temperatures.

50


H (kOe)

Figure 7. Relative change in thermal resistance of copper as a function of

magnetic field at several temperatures.

53
I nconel

718

T . T T

Figure 8. Characteristic curves of p q (zero field electrical resistivity),

Ap/Po (relative electrical resistivity change in a magnetic field),

and ^W/Wq (relative thermal resistivity change in a magnetic

field) for Inconel and copper below 20 K.

54
xlO"

14 . 0

0 kOe

5 10 15 20

T (K)

Figure 9. Lorenz ratios of Inconel 718 as a function of temperature at

several magnet fields.

55
56
NBSIR

SEMI-ANNUAL REPORT ON MATERIALS RESEARCH

IN SUPPORT OF SUPERCONDUCTING MACHINERY

FATIGUE AND FRACTURE TOUGHNESS

TESTING AT CRYOGENIC TEMPERATURES

R. L. Tobler, R. P. Mikesell, R. L. Durcholz, and R. P. Reed

Cryogenics Division

Institute for Basic Standards

National Bureau of Standards

Boulder, Colorado 80302

October 1974
Stjmmary: Fatigue and Fracture Toughness Testing at Cryogenic Temperatures

Fatigue crack growth rate (da/dN) , fracture toughness (K , K (J)) and

tensile properties have been determined at 298, 76, and 4 K for the following

structural alloys: Ti-5Al-2 . 5Sn ; Ti-6A1-4V; AISI 304, 310, 316; Inconel 750,

and A- 286. The titanium alloys Ti-5Al-2.5Sn and Ti-6A1-4V exhibited similar

properties: the fatigue crack growth rates were temperature independent over

the interval 298 - 4 K. Their fracture toughness at 4 K are equal to within

1/2 3-4 ksi • in . The stable austenitic alloys AISI 310, A-286, and Inconel

750 exhibit lower crack growth rates at 76 K and 4 K than at 298 K. The

fracture toughness of solution treated and aged A-286 alloy decreases by

only 10% as temperature decreases from 298 - 4 K; the fracture toughness

1/2

at 4 K is 109 ksi • in . The fracture toughness of solution treated and

double aged Inconel 750 appears to decrease by about 10% over the temperature

1/2 interval 298 - 76 K. The fracture toughness at 76 K is 173 ksi • in

AISI 310 stainless steel has superior resistance to fatigue crack growth at

4 K, compared with types 304 and 316 stainless steels. Conversely, AISI 304

and 316 stainless steels exhibit higher fracture toughness. These differences

appear to be related to the relative stability of the austenitic structure.

58
Contents : Fatigue and Fracture Toughness

Page

1. Introduction 60

1.1 Linear Elastic Fracture Mechanics 60

1.2 J-Integral Analysis 61

2. Experimental 62

2.1 Materials 62

2.2 Methods 63

2.2.1 Tensile Tests 63

2.2.2 Low-Temperature Fatigue 63

2.2.3 Fracture Toughness 66

3. Results 67

3.1 Tensile Properties 67

3.2 Fatigue Crack Growth Rates 68

3.2.1 Ti-5Al-2.5Sn 68

3.2.2 A-286 69

3.2.3 Inconel 750 69

3.2.4 AISI 310 69

3.3 Fracture Toughness 69

3.3.1 Ti-5Al-2.5Sn 69

3.3.2 A-286 70

3.3.3 Inconel 750 72

3.3.4 AISI 310 72

4. Discussion 73

5. References 77

List of Tables 82

List of Figures 83

59
FATIGUE AND FRACTURE TOUGHNESS TESTING AT CRYOGENIC TEMPERATURES

R. L. Tobler, R. P. Mikesell, R. L. Durcholz and R. P. Reed

1. Introduction

Recent Interest in superconducting machinery has prompted research on both the physical and mechanical properties of materials at A K. Since most

structural components of superconducting machinery are cooled with liquid

helium to 4 K, there is a need for mechanical property data in this environment.

Conventional tensile and impact tests do not provide design data for

predicting the load carrying capability of structures containing fatigue cracks

or other common flaws. The safe operating lifetime of a structure and the

possibility of brittle failure can be predicted only by fracture mechanics

analysis. The essential design parameters are the fracture toughness

(Kj^ or J-j-q) of a material and the fatigue crack growth rate (da/dN).

Some highly stressed components of superconducting motors and generators

are expected to sustain billions of fatigue cycles in their lifetime. Prototype machinery were constructed using some very tough austenitic alloys (300-

series stainless steels, Inconel 750) and elaborate fabrication techniques

employed to lower the probability of failures. Fracture mechanics data at

4 K, previously non-existent, are needed to enable judicious selection of

structural materials and accurate lifetime predictions for critical machinery

components.

This report presents fatigue crack growth rate and fracture toughness

data at 298, 76 and 4 K for the alloys Ti-5Al-2. 5Sn, A-286, Inconel 750,

and AISI 310 stainless steel. These data are compared to the results obtained

in the first six months on AISI 304, AISI 316, and Ti-6A1-4V.

1.1 Linear Elastic Fracture Mechanics

When a negligible amount of plastic deformation occurs prior to fracture,

the ASTM Standard Method of Test for Plane Strain Fracture Toughness of Metallic

Materials (ASTM E- 399-74) [1] can be applied. This method is based on the theory

of linear elastic fracture mechanics as discussed in references [2-4]. The

approach utilizes the stress intensity factor, K, to describe fracture behavior.

60
For test specimens of standard geometry, K is calculated from the relation:

K =

P

(1)

BW,1/2

where P is applied load, B is specimen thickness, W is specimen width, a is

crack length, and f(a/W) is a factor dependent on relative crack length.

A normally ductile material may fail in a relatively brittle manner in

the presence of a fatigue crack and under conditions of plane strain. Linear

elastic plane strain conditions prevail for thick specimens where the plastic

zone at the crack tip is constrained by the surrounding bulk of elastically

loaded material. Under these conditions, materials loaded in tension exhibit

a critical stress intensity, K^^, at which failure occurs catastrophically

without significant plastic deformation.

The parameter K^^ is a material property and a useful design criterion.

Provided that specimen size requirements are met, Kj.^ may be calculated from

eq (1) using the load determined from the fracture test record according to

the ASTM E-399-74 method. The size criterion for assuring linear elastic

behavior and valid K^.- data is that:

where is the 0.2% offset yield stress of the material. Subsized specimens

result in invalid fracture toughness measurements; these data are referred to

1.2 J- Integral Analysis

Specimens that do not satisfy eq (2) exhibit nonlinear elastic fracture

behavior. Plastic deformation at the crack tip becomes appreciable for thin

specimens and the concept of stress intensity loses its significance. In such

a case, the J-integral fracture criterion may be applied.

The J-integral approach is based on the premise that fracture behavior

is governed locally by crack tip conditions, even if plastic deformation precedes

fracture. J is an energy term derived from the area under the load versus deflection curve of a precracked specimen loaded in tension or bending. The

TC

(2)

as KQ*

61
critical value, Jj„, is the value of J just prior to crack extension—the

value required to initiate fracture.

The theory of the J- integral was originated by Rice [5]. Subsequent

experiments [6-10] verified that J is a useful fracture criterion that may

be applied in cases where the specimen exhibits linear elastic, elasticplastic, or fully plastic load-deflection behavior.

Although the J- integral criterion has been successfully applied to cases

of large-scale plasticity, a size criterion is necessairy to insure that J- integral

test results will be independent of specimen geometry and dimensions. A tentative size criterion for the J- integral test is that [10]:

B,b,a, > 50 (J/Ofio^ (3)

where b is the ligament of uncracked specimen and ^ ^-^^^ is the average of the

0.2% offset yield stress and the ultimate flow stress. A very practical con- sideration is that B is typically an order of magnitude smaller than the value

required for linear elastic tests as calculated from eq (2).

J^^ values can be converted to K^^ values using the relation [6, 7, 10]:

1/2

^IC = * '

where E is Young's modulus and v is Poisson's ratio. Thus it is possible to

derive K values from J-integral tests on smaller specimens than would normally

be required for linear elastic K tests. In this paper, the symbol K (J) is

used to distinguish values obtained using eq (4) from K data determined according to ASTM E- 39 9- 74.

2. Experimental

2.1 Materials

The materials tested in this study were the alloys Ti-5Al-2. 5Sn, A-286,

Inconel 750, and AISI 310 stainless steel. These materials were obtained in

the various forms listed in table 1. The Ti-5Al-2.5Sn and AISI 310 alloys

were tested in the as-received condition. Specimens of alloys A-286 and

Inconel 750 were machined from stock and subsequently heat treated prior to

testing.

62
A-286 was solution treated (1172 K (1650 F) , 2 hrs, oil quench) and

aged (1005 K (1350 F), 16 hrs, air cool). Inconel 750 was solution treated

(1255 K (1800 F), 1 hr, air cool, and double aged (1144 K (1600 F) , 24 hr,

air cool, followed by 977 K (1300 F) , 20 hr, air cool).

The chemical analyses of all materials are given in table 2. Note that

the Ti-5Al-2.5Sn alloy contained normal levels of the interstitial elements

C, N, 0, and H; it was not an extra- low- interstitial grade.

2.2 Methods

2.2.1 Tensile Tests

Smooth-bar uniaxial tensile tests were conducted to determine the 0.2%

offset yield stress and ultimate tensile strength for each material. The

tests were conducted according to the ASTM Standard Methods of Tension Testing

of Metallic Materials (ASTM designation E8-69) fll].

Testing was performed with a 10,000 lb (44.5 kN) capacity machine that was equipped with a titanium-stainless steel cryostat. The cryostat was

described by Reed [12]. The tests were performed using a crosshead rate of

-4 -1 -3 -1

3.3 X 10 in-sec (8.4 x 10 mm-sec ).

Testing in liquid nitrogen (76 K) was accomplished by immersing the

specimen and cryostat in a single metal dewar containing the cryogen. Testing in liquid helium (4 K) was accomplished with a double glass dewar arrangement; the outer dewar contained liquid nitrogen.

Tensile specimens had the geometry illustrated in figure 1. All materials

were tested in the transverse orientation; the tensile axis was perpendicular

to the rolling direction or principal forging axis. Thus, the fracture plane

orientation of tensile specimens was equivalent to that of fracture toughness

specimens

Load was monitored with a 10,000 lb (44.5 kN) commercial load cell.

Specimen extension was measured with a commercial clip-on extensometer that was calibrated with a dial micrometer at each temperature. The sensitivity

-4

of the extensometer at 4 K corresponded to a strain increment of 3.65 x 10

per millimeter of chart paper.

2.2.2 Low Temperature Fatigue

Fatigue tests were conducted with a 22,500 lb (100 kN) capacity servo- hydraulic mechanical testing machine. This was a cryogenic model; the fatigue

63
actuator was positioned above the specimen level to accommodate the cryostat

and to permit access from below for assembling dewars. The testing machine

and associated equipment are shown in figure 2.

Tests were conducted using a 22,500 (100 kN) capacity cryostat. The

cryostat and dewar arrangement for tests in liquid helium are illustrated

in figure 3. A more detailed description of the design of this cryostat and

associated apparatus was given previously fl3].

Compact tensile specimens having the geometries shown in figures A and 5

were employed in all fatigue and fracture tests. All specimens were 1.50 inch

(3.81 cm) thick and were proportioned in accordance with ASTM standard E-399-74.

AISI 310 alloy specimens were machined with the crack starter notch parallel

to the rolling direction of the original stock — a TL orientation [1], For

all other materials the plane of the notch was parallel to the principal forging

axis — a TS orientation [1].

Fatigue tests were performed using load control. Dynamic fatigue loads

were measured and controlled to within 1 percent by means of a digital peak

recording device. The load cycle was sinusoidal at a constant frequency of

20 Hz and a constant stress ratio (minimum load/maximum load) of 0.1.

Most cract growth data were obtained during precracking of specimens for

fracture tests. For these specimens the final crack length was limited to a ratio a/W between 0.45 and 0.60. Fatigue loads and stress intensities were

limited to low or intermediate values as shown in table 3. A few specimens

were tested solely for crack growth data at high stress intensities; for these

specimens the cracks were extended to a/W = 0.70 using loads up to 10,500 lbs

(47 kN). . ASTM type clip gages were used at all testing temperatures. Commercially

available foil resistance strain gage films were bonded to the clip gage beams

using an adhesive recommended for low temperature applications. The clip

gages were calibrated at temperatures from 298 K to 4 K, but sensitivity was not significantly temperature dependent. Clip gage linearity met ASTM standards

for K testing [1].

For linear elastic K tests of Ti-5Al-2. 5Sn, deflections were measured

J- v.* between attachable knife edges located at the specimen edge as shown in figure

4. For all other test specimens a modification in notch configuration was introduced to measure load-line deflections as shown in figure 5. Roberts' [14]

K calibration was used in each case.

64
Crack growth was monitored by compliance measurements. Direct measurement of crack lengths at the specimen surface was not possible during

cryogenic tests. The compliance technique Is based on the fact that for a given load, the specimen compliance (reciprocal of stiffness) Increases with

Increasing crack length. Compliance methods are discussed In detail elsewhere [4, 15].

Experimental crack- length — compliance correlations were obtained for

each material and temperature. An X-Y recorder was used to measure the com- pliance for a number of specimens of varying crack lengths. Fatigue strlatlons

were generated by load changes during precracklng and crack growth rate tests

so that several crack- length — compliance measurements were obtained from

each specimen. The data were plotted as a function of crack length, which

was measured as an average of three readings at the center and quarter points

of specimen thickness.

Theoretical crack-length — compliance curves were calculated from the

generalized solution given by Roberts [14], but these were not In complete

agreement with the experimental correlations. Comparison showed that the

calculated curves consistently underestimated the true crack lengths. The

difference between the calculated and experimental curves usually amounted

to about 5%.

The procedure for determining crack growth rates Involved plotting the

static compliance at Intervals during the tests. Using the experimental

crack-length — compliance correlation, crack length could then be obtained,

within ± 2 percent. The crack length, a, was plotted versus the total number,

N, of fatigue cycles, and the crack growth rate (da/dN) was determined by

graphical differentiation of the a-versus-N curve.

It is well known that the rate of crack growth can be retarded by changing

from a higher to a lower stress intensity range. Crack growth rates can also

be influenced by abrupt temperature changes. Whenever these kinds of changes

were made the crack growth rate was allowed to stabilize before data were

accepted as valid. Crack growth data under the new conditions were ignored

until the crack had advanced an increment equal to or greater than the plastic

zone size as calculated according to Irwin's formula fl6].

65
2.2,3 Fracture Toughness

Fracture tests of Ti-5Al-2. 5Sn were performed with the 22,500 lb.

(100 kN) servo-hydraulic test machine. The tests were performed using

-3 -1 -1

stroke control at a rate of 3.9 x 1- in • sec (100 mm • sec ).

J- integral fracture tests of A-286, Inconel 750, and AISI 310 alloys

were performed using a 60,000 lb (267 kN) capacity hydraulic test machine.

K was measured using the same J-integral test specimens. The tests were

-4 -1 -3 -1

conducted at a crosshead rate of 3.3 x 10 in • sec (8.4 x 10 mm • sec ) The 60,000 lb (267 kN) capacity cryostat and associated fracture testing

apparatus were described previously [13].

The majority of fracture tests were conducted at temperatures of 298,

76, and 4 K. A single test of Ti-5Al-2.5Sn was performed at 173 K in an atmosphere of nitrogen vapor. The specimens were allowed twenty to sixty

minutes to reach thermal equilibrium before testing.

A temperature-control system with a chromel-constantan thermocouple

was used to maintain a temperature of 173 K ± 4 K. The control thermocouple

was located on the specimen at the level of the fatigue crack. Vapor from

a pressurized liquid nitrogen dewar was introduced to the cryostat through

a solenoid valve. The valve admitted cold vapor to the cryostat as required

by opening and closing automatically in response to the electrical signal

from a commercial temperature recorder.

The specimens were precracked at the temperature of subsequent fracture

testing unless otherwise noted in the text. Ti-5Al-2.5Sn specimens were

precracked according to ASTM E-399-74 criteria. The precracking conditions

for other alloys are listed in table 3.

For the compact tensile specimens described in this report, J was calculated from the equation [17]

where A is the area under the load-deflection curve to a particular value of

displacement. The critical value, J-j-^j was taken as the value of J just prio:

to material separation. For elastic-plastic or fully plastic fracture tests.

66
the Jj.^ value was obtained by extrapolation of the J versus Aa curve. The extrapolation technique was based on a method outlined by Landes

and Begley flO]. Following their procedure, 3 to 5 nearly identical specimens with average crack lengths agreeing to approximately 3 percent were tested

at each temperature. A series of load-deflection curves were obtained by loading each specimen to a level sufficient to cause decreasing amounts of crack

extension. The specimens were then mloaded and heat tinted''" to oxidize the

surfaces where crack extension had occurred. The specimens were subsequently

pulled apart at either 298 K or 76 K. The average values of Aa were measured

and plotted versus the corresponding values of J.

In some tests a measurable deformation at the crack tip preceded material

separation. Apparent crack extension therefore included a contribution due

to deformation in addition to actual material separation.

To account for apparent crack extension due to deformation, the J/2a^^^^

line was plotted as described by Landes and Begley [10]. The intersection of

the J versus Aa curve and the J/2cj-, line represents the initial point of

flow

material separation; J was taken as the value of J at the point of intersection.

For AISI 310 specimens the length of the stretch zone was also measured

directly from the fracture surfaces with a traveling microscope. An average

value from several specimens was plotted as a vertical line on the J-Aa curves. When the J-Aa graphs appeared to be nonlinear beyond the stretch zone region, a concave upward curve was used to extrapolate the data. The concave

curve results from increasing over estimation in J with increasing crack ex- tension [10]. In some cases a well defined curve in the J-Aa plots was not

exhibited due to data scatter. Uncertainty in the value of J^^ was accounted

for by examining a number of reasonable extrapolations.

Kj^(J) values were calculated from J^^ values using eq (4). Data for

E and V at the temperatures of interest were taken from the work of Ledbetter,

Naimon, and Weston [18J.

3. Results

3.1 Tensile Properties

Tables 4-7 list the yield and tensile strengths for all alloys. The

results are shown as a function of temperature in figures 6-8. The stress1 Specimens were heat tinted for 15-30 minutes at 1300 - 1400 F.

67
strain curves are given in figures 9-12.

The yield strength and tensile strength of the Ti-5Al-2.5Sn alloy

nearly doubled with a decrease in temperature from 298 K to 4 K. At a value of 225 ksi (155 x 10^ Nm~^) the yield strength in liquid helium

exceeds that of the other alloys tested. However, ductility is low at 4 K, as indicated by the stress-strain curves of figure 9.

The tensile behavior of AISI 310 was similar to that reported previously

[13] for AISI 304 and 316 stainless steels. Typically, these alloys exhibit

high ductility, work hardening capability at temperatures as low as 4 K, and

discontinuous yielding behavior at 4 K. As shown in figures 10-12 the AISI

310, A-286, and Inconel 750 alloys plastically deformed via the discontinuous

yielding process at 4 K. This was observed for AISI 304 and 316 alloys also [13].

Figure 8 shows that the yield strength of AISI 310 increases considerably

with decreasing temperature over the interval 298 - 4 K. However, the yield

strength at 4 K was lower in comparison to the values obtained for A-286, Inconel

750, and Ti-5Al-2. 5Sn.

3.2 Fatigue Crack Growth Rates

3.2.1 Ti-5Al-2.5Sn

The preliminary room temperature crack growth rate data previously re- ported [13] for the Ti-5Al-2.5Sn alloy were inaccurate and have been corrected.

The preliminary data were calculated using Roberts' [14] theoretical compliance

correlation. The data have been corrected using a better compliance correlation, experimentally determined.

Fatigue crack growth rate data for Ti-5Al-2.5Sn are shown in figure 13.

The data represent tests of at least three specimens per temperature. At

298, 76, and 4 K, the rates are nearly equivalent. Any effect of temperature

is indistinguishable compared to the degree of scatter among specimens. All

crack growth rate data fit within a scatter band having a width of about

1/2 -2 1/2

5 ksi * in (5.5 MNm • m ). The scatter in data is somewhat greater

at 4 K than at 298 or 76 K.

Also shown in figure 13 is a straight line approximation of the data previously reported [13] for Ti-6A1-4V. The crack growth resistance of the

Ti-5Al-2.5Sn alloy is superior over the range of stress intensities examined.

68
3.2.2 A-286

Fatigue crack growth rate data for the A-286 alloy are shown in figure 14.

Most of this data was reported earlier [13].

Results at 298, 76, and 4 K demonstrate that cracks propagate at lower

1/2 rates at cryogenic temperatures than at room temperature. At ZiK = 40 ksi • in

-2 1/2 (43.9 MNm • m ), the crack growth rate at 76 K or 4 K is lowered by a factor

of two from the room temperature value. There is no measurable difference in

crack growth rates at 76 and 4 K.

3.2.3 Inconel 750

Fatigue crack growth rate data for Inconel 750 are shown in figure 15.

The scatter in data at each temperature was particularly low.

The crack growth rate decreased with decreasing temperature. The rate of

decrease between room temperature and 76 K was large and similar to that observed

for A-286. Similarly, there was only a small reduction in growth rate between

76 K and 4 K.

3.2.4 AISI 310

The results for AISI 310 stainless steel are shown in figure 16. Similar

to A-286 and Inconel 750, the crack growth rate is lower at cryogenic temperatures than at room temperature. A relatively large reduction in crack growth

rate occurs between 298 K and 76 K, but there appears to be no measurable

difference between 76 K and 4 K. In Figure 17 the results at 4 K for 310 stainless steel are compared

to results previously obtained [13] for types 304 and 316. Wide bands of

scatter were observed in the results for types 304 and 316. Over the range

of crack growth rate from about da/dN = 10 ^ - 10 ^ in • cycle (2.54 x 10

-4 -1

2.54 X 10 cm • cycle ), figure 17 indicates that cracks propagate more

rapidly for AISI 304 and 316 stainless steels than for 310. The crack growth

rate for 310 stainless steel is lower by a factor of about two.

3.3 Fracture Toughness

3.3.1 Ti-5Al-2.5Sn

Fracture test records for Ti-5Al-2. 5Sn specimens are illustrated in

figure 18. At cryogenic temperatures the load-deflection curves were ideally

linear elastic; failure occurred instantaneously at the critical stress intensity. A minor deviation from linearity was observed only in tests at room

temperature, where pop-in (unstable crack extension) occurred. Each pop-in

was marked by a discontinuity in the terminal portion of the test record just

prior to failure.

69
Table 8 lists the values of and K^^ for ten specimens. The specimen

size criterion was amply satisfied, as demonstrated in table 9.

Several results were invalid according to a strict application of

ASTM E-399-74. The data for three specimens were designated because of

minor deviations from fatigue crack uniformity. The invalidating criteria are

cited in table 9. However, these irregularities appear to have had no influence; scatter in the data is quite low, and values are in good agreement

with values at each temperature.

Figure 19 illustrates the temperature dependence of fracture toughness

for the Ti-5Al-2.5Sn alloy. Other results [22-23] for extra-low-interstitial

(ELI) Ti-5Al-2.5Sn alloys are presented, as well as data previously reported

for a mill annealed Ti-6A1-4V alloy [13] . The room temperature fracture tough- ness of Ti-5Al-2.5Sn is considerably higher than that of the Ti-6A1-4V alloy,

but there is only a small difference in K values at A K. A marked temperature dependence reduces the fracture toughness of the Ti-5Al-2.5Sn alloy from

1/2 -2 1/2 1/2 -2

70.3 ksi • in^ (77.2 MNm • m^ ) at 298 K to 38.3 ksi • in^ (43 MNm • 1/2

m ) at 76 K or 4 K, a decrease of 46%.

The fracture surface appearances of specimens tested at room temperature

and in liquid helium were predominantly flat. The width of the shear portion

at the specimen edge is less than 3% of total thickness. The shear portion

is practically non-existent at 76 or 4 K.

In load-deflection behavior, in K values, and in fracture surface

appearance, the results at 76 K and 4 K were indistinguishable.

3.3.2 A-286

Fracture test records for A-286 specimens are shown in figure 20.

Deviations from linearity were noticeable in the load- deflection curves at

each temperature. The degree of linearity increased at lower temperatures

but the ASTM E-399-74 specimen thickness criterion for linear-elastic behavior

was never satisfied. Discontinuities in load- deflection behavior and audible

pop- in phenomena occurred at all temperatures. Heat-tinting experiments re- vealed that small increments of stable crack extension always preceded pop-in.

The amount of stable crack extension that could occur prior to pop-in was estimated to decrease from about 0.07 inch (.178 cm) at 298 K to about 0.01

inch (0.025 cm) at 4 K.

70
J- integral test results are given in figure 21 and in table 10. As

illustrated in figure 21, the J/2a^, lines at 298, 76, and 4 K are rather

r low

steep and coincide. The points of intersection of the J-Aa curves with the

-3 -3

J/20^^^^ lines occurs for each temperature at Aa 10 inch (2.54 x 10 cm).

Thus, the deformation at the crack tip prior to material separation is quite

small; a stretch zone was not visible on the fractured surfaces of test

specimens.

At 76 and 4 K the J^^ values are well defined. In one test at 76 K, an average crack extension of only 0.004 inch (0.01 cm) was observed. The

heat-tinted fracture surface of this specimen revealed that crack extension

initiated d is continuously at several locations along the crack front. The

value of J for this test was very nearly the critical value, J-^p.

The scatter in the data at 298 K and the lack of data at very low values

of Aa contributed to uncertainty in the room temperature results. A number

of extrapolations indicate that lies in the region from 420 - 490

-2 -2

in • lb • in (0.073 - 0.086 MJ • m ).

The values shown in table 11 are not valid K^^ data because the

linear-elastic specimen thickness requirement was not satisfied. The data

of table 12 indicate that a 1.75 inch (4.44 cm) thick specimen is required

to assure valid K results at 4 K. The required thickness is significantly

greater at higher temperatures.

Shown in table 11 are calculations of the specimen strength ratio"^,

Rg^ fl]) at 298, 76, and 4 K. The and K-j.^(J) parameters are shown as a function of temperature in figure 22. Of these parameters, only the Kj^(J)

values represent a meaningful measure of fracture toughness. The R^^ values

illustrate qualitatively the effect of temperature on fracture toughness.

The J-integral results demonstrate that the A-286 alloy undergoes only

a small reduction in fracture toughness over the temperature interval 298 - 4 K;

K (J) decreases from 119 ksi • in"'"''^ (130.6 MNm"^ • m"*"^^) at 298 K to 109

1/2 -2 1/2 ksi • in (119.7 MNm • m ) at 4 K, a decline of about 10%.

'2P (2W+a)

1 The specimen strength ratio, R = r , is the ratio of the

B(W-a) a maximum nominal net-section stress to the tensile yield strength.

71
3.3.3 Inconel 750

Fracture data for Inconel 750 at 4 K have not been completely evaluated.

The results obtained at 298 K and 76 K are reported here.

Fracture test records for Inconel 750 are shown in figure 23. At 298 K

and 76 K fracture occurred by stable, ductile tearing. The load-deflection

curves deviated substantially from linearity. Valid K data could not be obtained according to the ASTM E-399-74 method due to inadequate specimen thick- ness as indicated in table 13. The values obtained by this method are listed

in table 14. The J- integral tests results are shown in table 15 and in figure

24. The size criterion for J- integral tests is amply satisfied as shown in

table 13.

At 298 K the J-Aa curve is concave upward. At 76 K the trend of the

J-Aa curve is obscured by scatter, but the J value is rather well defined

because two data points are very near the J/2df^- line. flow

Figure 25 indicates that the fracture toughness of Inconel 750 decreases

from a K (J) value of 189 ksi • in"*"^^ (207.5 Nm~^ • m"*"^^) at 298 K to a 1/2 -2 1/2 value of 173 ksi • in ' (190 Nm • m ' ) at 76 K, a decrease of 10%.

3.3.4 AISI 310

Fracture tests of AISI 310 at 4 K are now in progress. Results at 298 K

and 76 K are presented here.

An outstanding characteristic of room temperature tests was the large

amount of plastic deformation that occurred at the crack tip prior to material

separation. Loading caused a noticeable blunting of the crack tip and an apparent crack extension of about 0.08 inch (0.2 cm). The deformation was visible on the surfaces of fractured specimens as a zone of stretch between

the fatigue cracked region and the torn ligament. This deformed zone curved

upward, out of the plane of the fatigue crack.

The fracture test records for AISI 310 are shown in figure 26. The loaddeflection curves Indicate fully plastic behavior. At 298 K and 76 K the

mode of fracture was stable, ductile tearing. Discontinuities occur in the

4 K curves probably due to adiabatic, localized specimen heating.

J-integral test results are shown in table 18 and in figure 27. At room

temperature the J/^'^f2.ow ^ slope nearly equal to the slope of the

extrapolated J-Aa curve. The vertical stretch zone line, the J/2a.^ line, flow

72
and the J-Aa curve appear to meet at a point of triple intersection. The

J/2a^^^^ line appears to be in agreement with the datum point for a specimen

tested to a Aa value in the stretch zone. At 76 K the ^/^o^^^^^ line predicts a larger stretch zone than was actually

measured. The value of J for this temperature was determined by extrapolating

the J-Aa curve, to the measured value for the stretch zone. The first three

data points were weighted more heavily; this yields a more conservative estimate

of than if the fourth datum point were equally weighted. The value given

for at 76 K is preliminaiy; additional data are needed at Aa < 0.015 inch

(0.38 cm).

The values listed in table 16 do not represent valid fracture toughness

measurements. The ratio of fracture toughness to yield stress is very high at

each temperature; hence, linear elastic plane strain behavior could never be

observed in practical test specimen sizes. Table 17 indicates that at 76 K

linear elastic behavior would be observed in 18 inch (46 cm) thick specimens.

On the other hand, specimens '^0.6 inch (1.5 cm) are sufficient for J- integral

tests.

The Kj^(J) values converted from the J^^ data are shown in figure 28.

The K^(|,(J) value at 76 K is substantially higher than the value at 298 K—

a

clear indication that the fracture toughness of AISI 310 increases at cryogenic

temperatures.

The fracture surfaces of AISI 310 specimens were not ferromagnetic, in

contrast to the findings reported for AISI 304 and AISI 316 stainless steels

[13]. This indicates that there was no austenite-to-martensite transformation

for the AISI 310 alloy, even when deformed and fractured in liquid helium.

This is in agreement with other observations that the austenitic 310 alloy

remains stable during deformation at low temperatures whereas 304 and 316

stainless steels do not [25, 26].

4. Discussion

The tensile results for the alloys Ti-5Al-2.5Sn, Inconel 750, and

AISI 310 were in close agreement with data from the literature [12, 21-23,

27-38] but the yield strengths determined for the precipitation hardened

A- 286 alloy were 15% lower than might be expected. Data in the literature

[37-45] indicate that room temperature yield strengths of about 105 ksi

7 -2

(72.4 X 10 Nm ) and higher can be achieved. Probably due to the complex

73
metallurgical structure of alloy A-286, It has been found [46] that the optimum

tensile properties are not always achieved with the standard heat treatment.

Fracture toughness data for Ti-5Al-2.5Sn at cryogenic temperatures

are available in the literature [22, 23, 47-51] but the existing data are incomplete. Some data [47, 50, 51] obtained for thin sheet materials

(B < .5 inch) were not sufficient to provide valid K data over the entire

cryogenic to ambient temperature range. No data at 4 K had been reported.

Although it is generally accepted that ELI grades of Ti-5Al-2.5Sn

exhibit superior low temperature toughness compared to normal interstitial

grades, there has been a lack of quantitative data relating K values to

J- V* interstitial content. The present results should provide a useful basis for

comparison.

In figure 19 the data of Pyle et al. [22] and Carman et al. [23] were

presented. These authors performed tests on ELI grades of Ti-5Al-2.5Sn

having a total Fe+C+C+O+N content of 0.28 - 0.46%. At 20 K, the

1/2

average K values are 9 ksi • in higher than the normal interstitial

J- \j (Fe + C + 0 + N = .52%) Ti-5Al-2.5Sn alloy reported here. It may be significant

to note that the ELI grades exhibited considerable variability. Scatter in

the results of Pyle et al. and Carman et al. amounted to ± 14%.

The K (J) values for Inconel 750 at 76 K were in the range 161-185

1/2 ksi • in . Scatter in the data between specimens may have been due partly

to the fact that the specimens were obtained from individual forgings. Data

reported by Lessman, Logsdon, Kossowsky, Mathur, and Wells [21] showed

1/2 Kj^(J) values of about 70 ksi • in for Inconel 750 at 76 K. The material

tested by Lessman, et al. was solution treated and double aged using different

heat treatment conditions than those applied in the present investigation.

Thus, processing and heat treatment apparently influence the fracture toughness

of Inconel 750 significantly.

Depending primarily on alloy composition, the AISI 300 series austenitic

stainless steels are metas table. At low temperatures the austenitic phase (y)

may transform to bcc (a' ) or to hep (e) martensitic structures. Literature

74
is available describing the effect of transformation on low temperature

strength and ductility [25, 52-55] but little has been published concerning the effect of such transformations on fatigue crack growth rates and

plane strain fracture toughness.

In the present studies, magnetic measurements revealed that cyclic

loading of AISI 304 and 316 stainless steels during crack growth experiments

at 4 K caused considerable formation of martensite. There was no evidence

of transformation for AISI 310, agreeing with tensile data [25, 26], As

much as 95% martensite may result from strain induced transformations in

AISI 304 and 316 alloys [25, 26].

When crack growth rate data for AISI 310 were compared with previous

results [13] for AISI 304 and 316 alloys, the crack growth rates at room temperature were nearly equivalent. At room temperature, the three alloys are essentially stable and the data represent the rate of crack propagation

through austenite.

During tests of AISI 304 and 316 alloys at 76 and 4 K, it can be assumed

that the martensitic transformation occurred locally in the plastic zone ahead of the crack tip; transformation preceded material separation. The

parameter measured in these tests was the rate of crack propagation through

partially martensitic material. But, the data for AISI 310 represent the rate

of propagation through stable austenite. The comparison in figure 17 for 4 K

data indicated that over the range of stress intensity examined, the crack

growth rate in the metastable alloys (AISI 304, 316) is higher. But, at room temperature all alloys (304, 310, 316) exhibit equivalent crack growth rates.

It is suggested that the crack growth rate behavior of AISI 300 series stainless steels is influenced by the martensitic transformation; apparently,

transformation preceding crack growth results in higher crack growth rates.

The AISI 310 alloy exhibits excellent fracture toughness at cryogenic

temperatures, but the fracture toughness of 304 and 316 stainless steels

are greater. The austenite to martensite transformation provides a plausible

explanat ion for the difference since the transformation is an energy absorbing

process. For materials that transform during fracturing, not all the energy

absorbed during the test goes into crack extension [56]. Martensitic phase

transformations account for the high fracture toughness observed in TRIP

(transformation induced plasticity) steels [57].

75
The mechanical properties of the materials that we have investigated

at 4 K are compared in table 19. The alloys are classified in three groups

that represent a wide range of mechanical behavior. The fracture toughness

data for these materials, based on the K and K (J) parameters, varies by

1/2 over an order of magnitude from 35 ksi • in for the Ti-6A1-4V alloy to

1/2 475 ksi • in for AISI 304. The same materials vary in yield strength

by a factor of three. The precipitation-hardened austenitic alloys occupy

an intermediate position, combining a balance of fracture toughness and

yield stress with high resistance to fatigue crack growth at cryogenic temperatures . The choice of materials for specific applications must be made on the

basis of several interrelated design criteria. The alloys in table 19 are ranked in descending order according to fracture toughness but on the basis

of yield strength the ranking would be exactly reversed. The fact that fracture

toughness and yield stress are inversely related is a typical example of the

trade-offs that are involved in materials selection.

76
5. References

1. Standard method of tests for plane strain fracture toughness of metallic

materials, E- 399-74, Annual Book of ASTM Standards, part 10, pp 432-451

(Amer. Soc. Test. Mater., Easton, Md. , 1974).

2. Fracture Toughness Testing and its Applications, ASTM STP 381, pp 1-409,

(Amer. Soc. Test. Mater., Baltimore, Md., 1965).

3. Brown, W. F., Jr., and Srawley, J. E., Plane Strain Crack Toughness

Testing of High Strength Metallic Materials, ASTM STP 410, pp 1-129

(Amer. Soc. for Test, and Mater., Baltimore, 1966).

4. Review of Developments in Plane Strain Fracture Toughness Testing,

ASTM STP 463, Ed. W. F. Brown, Jr., pp 1-170, (Amer. Soc. Test. Mater,

and NASA, Lutherville-Timonium, Md, 1970)

.

5. Rice, J. R. , A path independent integral and the approximate analysis

of strain concentration by notches and cracks, J. Appl. Mechs . , Trans.

ASME 35 , 379-386 (1968)

.

6. Begley, J. A. and Landes, J. D. , The J- integral as a fracture criterion,

in Fracture Toughness, Proc. of the 1971 National Symposium on Fracture

Mechanics, Part II, ASTM STP 514, pp. 1-20 (Amer. Soc. Test. Mater.,

Philadelphia, Pa., 1972).

7. Landes, J. D., and Begley, J. A., The effect of specimen geometry on J^^

in Fracture Toughness, Proc. of the 1971 National Symposium on Fracture

Mechanics, Part II, ASTM STP 514, pp. 24-39 (Amer. Soc. Test. Mater.,

Philadelphia, Pa., 1972).

8. Yoder, G. R. and Griffis, C. A., J- integral and the initiation of crack

extension in a titanitim alloy, NRL report 7662, Naval Res. Lab., (1973),

20 pp.

9. Kobayashi, A. S., Chiu, S. T. , and Beeuwkes, R. , A numerical and experimental investigation on the use of the J- integral. Eng. Fract. Mech. 5^, 298-305 (1973).

10. Landes, J. D. , and Begley, J. A., Test results from J- integral studies:

an attempt to establish a J testing procedure, in Fracture Analysis,

ASTM STP 560 (Amer. Soc. Test. Mater., Easton, Md., 1974).

11. Standard methods of tension testing of metallic materials, E8-69

,

Annual Book of ASTM Standards, part 10, pp. 90-110 (Amer. Soc. Test.

Mater., Easton, Md. , 1974).

12. Reed, R. P. , A cryostat for tensile tests in the temperature range

300" to 4*K, paper K-3, in Advances in Cryogenic Engineering, 7^, Ed. K. D. Timmerhaus, pp. 448-454 (Plenum, New York, 1962).

77
13. Tobler, R. L., Mlkesell, R. P., Durcholz, R. L., Fowlkes, C. W.

,

and Reed, R. P., Fatigue and fracture toughness testing at cryogenic

temperatures, in Semi-Annual Report on Materials Research in Support

of Superconducting Machinery, NBSIR 74-359, Nat. Bur. Stands.

(1974) 308 pp; available from NTIS, AD 780-596/3WM.

14. Roberts, E., Jr., Elastic crack-edge displacements for the compact

tension specimen. Mater. Res. Stand. 9^, 27 (1969).

15. Bubsey, R. T., Fisher, D. M. , Jones, M. H., and Strawley, J. E.,

Compliance measurements. Chap. 4, in Experimental Techniques in

Fracture Mechanics, Ed. A. S. Kobayashi, First Edition, pp. 76-95,

(Soc. Exper. Stress Anal, and Iowa State Univ., Westport, Conn., 1973).

16. Irwin, G. R., Fracture, Handbuch der Physick, 6^, pp. 551-590 (Springer

Verlag, Berlin, 1958).

17. Rice, J. R., Paris, P. C, and Merkle, J. G., Some further results on J-integral and estimates, in Progress in Flaw Growth and Fracture

Toughness Testing, ASTM STP 536, pp. 231-244, (Amer. Soc. Test. Mater.,

Baltimore, Md., 1973).

18. Ledbetter, H. M. , Naimon, E. R. , and Weston, W. F., Elastic properties

of engineering alloys at cryogenic temperatures, in Semi-Annual Report

on Materials Research in Support of Superconducting Machinery, NBSIR

74-359, Nat. Bur. Stand (1974), 308 pp; available from NTIS, AD 780-596/3WM.

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

Selected Materials from 20 to 300°K, Nat. Bur. Stand. Monograph 63,

pp. 1-51 (Washington, D.C., 1963).

20. Basinski, Z. S., The instability of plastic flow of metals at very

low temperatures, Proc. Roy. Soc. Lond. A240, 229-242 (1957).

21. Lessman, G. G., Logsdon, W. A., Kossowski, R., Mathur, M. P., and

Wells, J. M. , Structural materials for cryogenic applications,

74-9D4-CRYMT-RI, (Westinghouse Research Labs., Pittsburgh, Pa.,

1974) 59 pp.

22. Pyle, R., Schillinger, D. E., and Carman, C. M. , Plane strain fracture

toughness and mechanical properties of 2219-T87 aluminum and 5Al-2.5Sn

ELI titanium alloy weldments and one inch thick 5Al-2.5Sn ELI titanium

plate, NASA CR-72154, Frankford Arsenal, (1968) 86 pp.

23. Carman, C. M. , Forney, J. W. , and Katlin, J. M., Plane strain fracture

toughness and mechanical properties of 5Al-2.5Sn ELI titanium at room

and cryogenic temperatures, NASA CR-54296, Frankford Arsenal (1966),

59 pp.

78
24. Freed, C. N., and Kraft, J. M. , Effect of side grooving on measurements

of plane strain fracture toughness, J. Mater., 1, No. 4, 770-789 (1966).

25. Gunter, C. J., and Reed, R. P., The effect of experimental variables

including the martensite transformation on the low temperature mechanical

properties of austenitic stainless steels, Trans. Quart., 5_5, No. 3,

399-419 (1962).

26. Larbalestier, D. C. , and King, H. W. , Austenitic stainless steels at

cryogenic temperatures: 1 - structural stability and magnetic properties.

Cryogenics 13, No. 3, 160-168 (1973).

27. Espey, G. B., Jones, M. H., and Brown, W. F., Jr., Sharp-edge-notch

tensile characteristics of several high strength titanium-sheet alloys

at room and cryogenic temperatures, in Low Temperature Properties of

High Strength Aircraft and Missile Materials, ASTM STP 287, pp. 74-96,

(Amer. Soc. Test Mater. , Baltimore, Md. , 1960).

28. Hoke, J. H. , Mabus, P. G. , and Goller, G. N., Mechanical properties

of the stainless steels at subzero temperatures. Met. Prog., 55_, 643 (1949).

29. Kropschot, R. H. , and Graham, W. F. , Mechanical properties of the

austenitic stainless steels at low temperatures, NBS 5009, Nat. Bur.

Stand. (1956) 6 pp.

30. Deslsto, T. S., and Carr, F. L. , Low Teiqjerature Mechanical Properties

of 300 series stainless steel and titanium, Watertown Arsenal Laboratories,

(1960) 15 pp AD 609909.

31. Schwartzberg, F. R. , Osgood, S. H., Keys, R. D. , and Kiefer, T. F.

,

Cryogenic materials data handbook, AFML-TDR-64-280, Martin Marietta

Corp. (1964).

32. Belton, J. H. , Godby, L. L. , Taft, B. L., Materials for use at liquid

hydrogen temperature, in Low Temperature Properties of High Strength

Aircraft and Missile Materials, ASTM STP 287, pp. 108-121, (Amer. Soc.

Test. Mater., Baltimore, Md., 1960).

33. Shogan, R. P. , Tensile properties of irradiated Inconel 718 and

Inconel X-750 at cryogenic temperatures, WANL-TME-1922, Westinghouse

Electric, (1969) 38 pp.

34. Watson, J. F. , and Christian, J. L. , Low temperature properties of

K-Monel, Inconel-X, Rene 41, Haynes 25 and Hastelloy B Sheet Alloys.

J. Basic Eng., 84, 265-277 (1962).

35. Weleff, W. , McQueen, H. S., and Emmons, W. F. , Cryogenic tensile

properties of selected aerospace materials, paper A-2, in Advances

in Cryogenic Engineering, 10, Ed. K. D. Timmerhaus, p. 14-25 (Plenum,

New York, 1965)

.

79
36. Lovoy, C. V., Low-temp mechanical properties of Inconel-X and its

weldments, IN-P and VE-M-62-5, Marshall Space Flight Center (1962),

34 pp.

37. Low temperature mechanical properties of various alloys, a compilation,

NASA-SP-5921(01), NASA, (1970), 24 pp.

38. Martin, H. L. , Miller, P. C, Imgram, A. G. , Campbell, J. E., Effects

of low temperatures on the mechanical properties of structural metals,

NASA-SP-5012(01) , NASA, (1968) 35 pp.

39. Roberts, D. A., Roach, D. B., Hall, A. M. , Physical and mechanical

properties of nine commercial precipitation hardenable stainless steels.

Defense Metals Information Center Report 112 (1959)

.

40. Montano, J. W. , Mechanical properties of high strength A-286 bolts

at cryogenic temperatures, IN-P and VE-M-64-1, Marshall Space Flight

Center (1964)

.

41. Slunder, C. J., Hoenie, A. F. , and Hall, A. M., Thermal and mechanical

treatments for precipitation hardenable stainless steels and their

effect on mechanical properties, NASA TM-X-53578, Marshall Space

Flight Center (1967).

42. Schwanbeck, C. A. , Effects of interim warming on tensile properties

of A-286 stainless steel, LAC ER-10008, Lockheed-Georgia Co. (1968)

20 pp.

43. Masteller, R. D., et al. , Properties of cryogenically worked materials,

NASA CR-72638 (N70-27114) , Martin Marietta Corp., (1970).

44. Montano, J. W. , An evaluation of the mechanical and stress corrosion

properties of cold worked A-286 alloy, NASA-TM-X-64569 , NASA, (1971).

45. Anonymous, Armco A-286, Product data bulletin SA-1, Armco Steel

Corp. (1966) 8 pp.

46. Amer. soc. met. committee on heat resisting alloys. Heat treating

of heat resisting alloys, in Metals Handbook, 2, Heat Treating, Cleaning

and Finishing, 8th Edition, pp. 257-268 (1964).

47. Witzell, W. E., Fracture mechanics - plane strain characteristics of

several materials, GDC-ERR-AN-1228, General D3mamics/Convair (1967)

35 pp. AD 854611L.

48. Witzell, W. E., Fracture data for materials at cryogenic temperatures,

AFML-TR-67-257, Air Force Materials Laboratory (1967) 73 pp.

49. McClaren, S. W. , and Foreman, C. R. , Cryogenic design data for materials

subjected to uniaxial and multiaxial stress field, AFML-TR-65-140,

LTV Aerospace Corp. (1965) 144 pp.

80
50. Tiffany, C, Fo, Lorenzj P. Mo, and Hall, L. Ro, Investigation of

plane-strain flaw growth in thick walled tanks, NASA CR-54837, The

Boeing Company (1966) 63 pp.

51. Sullivan, T. L., Uniaxial and biaxial fracture toughness of extra- lowinterstitial 5Al-2.5Sn titanium alloy sheet at 20°K, NASA-TN-I>-4016

,

Lewis Research Center (1967) 20 pp.

52. Reed, R. P. , The spontaneous martensitic transformation in 18% Cr,

8% Ni steels. Acta Met., 10, p. 865-877 (1962).

53. Watson, J. F. , Christian, J. L. , Low temperature properties of coldrolled AI3I types 301, 302, 304 ELC, and 310 stainless steel sheet,

in Low Temperature Properties of High Strength Aircraft and Missile

Materials, ASTM STP 287, pp. 170-194 (American Soc. Test Mater.,

Baltimore, Md., 1960).

54. Watson, J. Fc , and Christian, J. L. , A study of austenite decomposition

at cryogenic temperatures. Trans. AIME, 224 , 998-1005 (1962).

55. Bhandarkarj Do, Zackay, V. F. , and Parker, E. R., Stability and mechanical

properties of some metastable austenitic steels. Met. Trans., 3^, 2619-2631

(1972).

56. Antolovich, S. D. , Fracture toughness and strain- induced phase transformations. Trans. AIME, 242, 2371-2373 (1968).

57. Gerberich, W. W., Hemmings, P. L. , Zackay, V. F. , Parker, E. R.

,

Interactions between crack growth and strain- induced transformation,

paper 24, in Fracture 1969, Ed. P. L. Pratt (Chapman and Hall Ltd.,

London, 1969) pp. 288-296.

81
List of Tables

Page

1. Material condition and form as received 84

2. Chemical analyses (wt %) 84

3. Fatigue cracking conditions 85

4. Yield and ultimate tensile strength for Ti-5Al-2.5Sn 86

5. Yield and ultimate tensile strength for A-286 86

6. Yield and ultimate tensile strength for Inconel 750 87

7. Yield and ultimate tensile strength for AISI 310 87

8. Fracture toughness of Ti-5Al-2.5Sn 88

9. Thickness required for valid K tests of Ti-5Al-2.5Sn .... 88

J- \^ 10. J-integral test results for alloy A-286 89

11. K and R parameters for A-286 89

12. Specimen thickness required for valid fracture tests of A-286 . 90

13. Specimen thickness required for valid fracture tests of

Inconel 750 90

14. parameters for Inconel 750 90

15. J-integral results for Inconel 750 91

16. parameters for AISI 310 92

17. Specimen thickness required for valid fracture tests of

AISI 310 92

18. J-integral results for AISI 310 93

19. Summary of mechanical properties at 4 K 93

82
List of Figures

Page

1. Tensile specimen (1 In = 2.54 cm) 94

2. Cryogenic fatigue testing apparatus 95

3. Liquid helium fatigue cryostat 96

4. Compact tensile specimen used for fracture tests of

Tl-5Al-2.5Sn; B = 1.50 In, W = 3.00 In (1 In - 2.54 cm) ... . 97

5. Compact tensile specimen used for J-lntegral tests of A-286

Inconel 750 and AISI 310 alloys (1 In - 2.54 cm) 98

6. Yield and tensile strengths of Ti-5Al-2.5Sn as a function of

temperature 99

7. Yield and tensile strengths of Inconel 750 as a function of

temperature 100

8. Yi°eld and tensile strengths of AISI 310 as a function of

temperature 101

9. Stress-strain curves for Tl-5Al-2.5Sn ....... 102

10. Stress-strain curves for A-286 103

11. Stress-strain curves for Inconel 750 104

12. Stress-strain curves for AISI 310 105

13. Crack growth data for Ti-5Al-2.5Sn 106

14. Crack growth data for A-286 107

15. Crack growth data for Inconel 750 108

16. Crack growth data for AISI 310 109

17. Crack growth data for AISI 310, 304 and 316 stainless

steels at 4 K 110

18. Fracture test records for Tl-5Al-2.5Sn (1 lb - 454 gm) .... Ill

19. Temperature dependence of fracture toughness of

titanium alloys 112

20. Fracture test records for A-286 113

21. J-integral versus crack extension for A-286 114

22. Temperature dependence of fracture toughness of A-286 115

23. Fracture test records for Inconel 750 (1 in - 2.54 cm,

1 lb = 4.47 N) 116

24. J-integral versus crack extension for Inconel 750 117

25. Temperature dependence of fracture toughness for Inconel 750 . 118

26. Fracture test records for AISI 310

(1 in - 2.54 cm, 1 lb - 4.47 N) 119

27. J-integral versus crack extension for AISI 310 120

28. Fracture toughness of AISI 310 at 298 K and 76 K 121

83
Table 1. Material condition and form as received.

Material Form Condition as Received

Ti-5Al-2. 5Sn

A-286

Inconel 750

AISI 310

4 3/4" square bar

4 1/2" square bar

4 1/4x4 1/4 X 1 3/4"

Forgings

1 . 5" Plate

Commercial annealed

Forged and annealed

Forged, hot-finished

Commercial annealed

NOTE: 1 inch = 2.54 cm.

Table 2. Chemical analyses (wt%).

Ti-5Al-2. 5Sn 5.28A1, 0. 327Fe, 2.46Sn, O.OIC, O.OIN, 0. 1740, 40 ppm H

AISI 310 Bal Fe, 0.

20.4Ni, 0.

038C, 1.41Mn, 0. 023P, 0.016S, 0. 72Si, 24.65Cr,

13Mo, 0.14CU

A -286 Bal Fe, 0.

24.97Ni, 1 05C, 1.

. 30Mo,

52Mn, 0.016P, 0.007S, 0.54Si, 13.96Cr,

2. 23Ti , 0. 004 B, 0.30 V, 0. 19 Al

Inconel 750 6.58Fe, 0.

0. 04 Cu, 2

04C, 0.

. 51 Ti,

20Mn, 0.007S, 0.25Si, 15.l6Cr, 73. 52Ni,

0. 78 Al, 0. 89 Cb + Ta

84
Table 3. Fatigue cracking conditions.

Alloy Temperatures a/W (P,/P

T TY\ f X 111 X Ti-5Al-2. 5Sn 298, 76, 4 K. .45 - .55 -0.5

f

. 6K^; K. ^ . Q f 002 E

A- 286 298, 76, 4 K ~ 0.55 0.4 < . 002 E

Inconel 750 298, 76, 4 K ~ 0. 60 0. 3 < 0014 E

AISI 310 298 K ~ 0. 63 0. 7 <.

• «Kq^

'^f ' 002 E

76, 4 K ~ 0. 63 0.4 K

f

<.

. 002 E

P /P is the ratio of maximum fatigue load/maximum load to fracture,

f m

K is the maximum stress intensity during fatigue, E is Young's modulus.

85
Table 4. Yield and ultimate tensile strengths for Ti-5Al-2. 5Sn

Temperature

(K)

0. 2% Yield Strength

(ksi)

Ultimate Tensile Strength

(ksi)

298 127.8

126. 0

Avg= 127

135. 2

133. 0

Avg= 134

225.4

N/A

Avg=225

226.4

224. 7

Avg= 226

NOTE: 1 ksi = 0. 689 x 10^ Nm"^

Table 5. Yield and ultimate tensile strengths for A -286

Temperature ^ . . 0. 2% Yield St rength Ultimate Tensile Strength Source ,t^\ Orientation /, . x . ^ (K) (ksi

)

(ksi)

This 298 Trans. 93. 3 150^

Report 83. 0

Avg = 88

76 Trans. 110.5 198^

106. 5

Avg = 108

4 Trans. 129.1 215^

128.7

Avg = 129

Warren 298 Long. Ill 160

and

Reed 195 Long. 120 176

76 Long. 135 209

20 Long. 150 • 235

NOTE: 1 ksi = . 689 x 10^ Nm"^

a An estimation based on the data of Warren and Reed [ 19] .

86
Table 6. Yield and ultimate tensile strength for Inconel 750.

Temperature 0. 2% Yield Strength Ultimate Tensile Strength

(K) (ksi) (ksi)

298 106.0 174.9

100. 7 166.

7

Avg=103 Avg=171

76 _ 119.6 205.9

115.

7

203.9

Avg=118 Avg=205

4 128.1 226.6

124.7 224.2

Avg= 126 Avg= 225

NOTE:l ksi = . 689 x lo"^ Nm"^

Table 7. Yield and tensile strength for AISI 310.

Temperature

(K)

0.2% Yield Stress

(ksi)

Ultimate Tensile Strength

(ksi)

298 35. 3

34. 2

Avg=35

78.9

78. 7

Avg=79

76 79.2

77. 7

Avg=78

151.7

153. 6

Avs=153

111.0 171

NOTEl ksi = 0. 689 x lo"^ Nm"^

Measured according to crosshead travel.

87
Table 8. Fracture toughness of Ti-5Al-2. 5Sn .

Temperature Specimen a/W Kj^(ksi-inl/2) KQ(ksi.in^/2)

1 C C 9 DO. 9

3 . 562 71. 7

. boi -7 1 ^.a 71. D

173 10 . 629 - 53. 2t>

r 0

QO

2 Q na -59. 9

9 .460 38. 3

12 .495 - 39.5^

4. A J 7, rt 13 .552 37.1

14 . 504 36.6^

^Fatigue crack plane deviated '^3° from notch plane; -10° is acceptable,

b.Fatigue stress intensity exceeded 0.60 Kq.

^Side grooved specimen; calculated according to Freed et al. [24].

Edge crack = 89. 5% of average crack length; ^90% is acceptable,

NOTE: 1 ksi • in^/^ = 1. 098 X lO^Nm'^-m^^^

Table 9. Specimen thickness required for valid K for Ti - 5A1-2. 5Sn ;

(K) (in) (in) (ksi) (ksi-inV^) (in)

Temperature a B a K 2.5(K /a„)

298 1, 35 - 1. 65 1. 5 122 70.3 0.835

76 1, 35 - 1. 65 1. 5 200 39. 2 0. 096

4 1. 35 - 1.65 1. 5 225 38,3 0, 072

NOTES: 1 in = 2. 54 cm.

1 ksi = , 689 X lO'^ Nm"^

1 ksi. inV2 ^ Q(^g ^ ^q6 j^fm"^ • m^^^
Table 10. J-integral test results for alloy A-286.

Temperature

(K) Specimen a/W

J

(in-lb-in-2)

L a

(in) (in* lb' in"

E/(l-v^)

) (lO^psi)

Kj^(J')

(ksi-in^/^)

7 Q Q (Co n 7 A . U f D 4

5

. 554

. 543

806

507

. 101

. 015

/ICC -1- Q07 J 1 . 1 1 Q -4- A C 07

1 c . Drt D DC / n "J Q

76 6 .542 396 . 005

8 . 556 658 . 05 6 385 ± 6% 33. 17 1 1 3 ± 3%

1 J . U X ^

4 0 . 550 402 . 035

7. • . 555 543 . 148 3 60 ± 4% 33. 22 109 ± 2%

10 . 548 373 . 022

NOTES: 1 in = 2.54 cm .

1 ksi • inl/2 ^ 1 098 X 10 Nm" • m

-2 -4 -2

1 in- lb- in =1.75x10 MJm

Table 11. K and R parameters for A-286

Q S C

Temperature

(K) Specimen i i. a/W

Kq

(ksi.inl/2) Ho

298 2 . 548 95

4 . 554 111

5 . 543 105 1.72

15 . 546 101

Avg=103

76 6 . 542 117

8 .556 135

1.49

13 . 545 122

9 . 565 122

4 0 . 550 120

7 . 555 114 1.20

10 . 548 116 jl Avg=117

NOTE: Iksi • in^^^ = 1. 098 • 10^ Nm"^ 1/2

• m

89
Table 12. Specimen thickness required for valid fracture tests of A-286

Temperature

(K)

2

2. 5 (K /a ) 2.

Q y

(in)

5 [K^(J)/a

(in)

2

y IC flow

(in)

298 3,1 4. 5 0. 20

76 3.3 2. 7 0.13

4 2.1 1.8 0.10

NOTE: 1 inch = 2. 54 cm

Table 13. Specimen thickness required for valid fracture tests of Inconel 750.

Temperature

(K)

2.5[K (J)/a ]^

IC y

(in) (in)

298 8. 6 .41

76 5. 7 .29

NOTE: 1 inch = 2. 54 cm

Table 14. K parameters for Inconel 750

Q

Temperature Specimen a/W

(K) (No. ) (ksi* in )

2 . 601 104

7 . 603 118

10 . 602 106

5 . 604 123

Avg =113

8 . 605 118

11 , 602 139

4 . 599 160

Avg = 139

6 2 NOTE: 1 = 1^ 098 x 10 Nm" • m

90
in

-H

CO

-Ho

o

00 00 O OO (\J o o o o

o o O O

o ro OO

fM O sD

1—1 fM 1—1 1—1

O O O ON

(\4 —( r-l

rg o in

00ON

-HCO

fM

OO

o

00

00

ON r-H O nD nO o

o o o o

o o o 00

CO 00 ON o fM

1—

1

1—

(

<M

ON ON O O

r-l 1—1 <M (M

QQ r-i ^

fM

in

fM

II

el

of—

4

in

fM

I

• r-l

h

fM

fM

nO

fM o1—

<

a

00oo

o1—1

>^ II ON fM 00sO 1—

1

a

II

.1-1

• 1-1 •1-1 (0 (0 (X

1—1

C/3 W

H

91
Table 16. Parameters for AISI 310

Temperature

(K)

Specimen

(No. ) a/W

K

(ksi? in 1/2)

298

r .

.4

2

9

5

. 628

. 630

. 635

. 640

40. 6

45.1

44.1

48, 0

Avg= 44. 5

1 o

113

8

. 618

. 618

. 609

80.8

77. 8

83. 2

A ^rt-* — Q A Q

1 ksi. in^/^ = 1. 098 X 10^ Nm-2. m^/^

Table 17. Specimen thickness required for valid fracture tests of AISI 310

Temperature

(K)

2.5[K (J)/ay]

( incli)

5° ( Jlc/"£low*

( inch)

298 54 1. 23

76 17. 6 0. 60

NOTE:l in = 2. 54 cm.

92
fM

I—I

CO \ o

I—I -H

T •I—I I oH

LD

Ln+ I

If) in+ I in

I—

r—I

1-4 CO + I Oo ^ un 00

rl^ o ^ m o o o o + 1

in

(N)

O 00 ^ Tt" o o o o o o S 3 in 5^ r—I I—I 00 O O o n£) i-H vo in CO

(NJ CO

M o

(M

CO (M o 00

(NJ rHO

pr vo

Tf< fM o in CO 00

00

CO I

00o

II

•r-i en

fM

It

in

00o

II in

II fl

•iH Si . o

.7i

X

WHO 93

COsH

a; Cl ^ .2

CO

CO

(0

t—4

0)

•rH oU

(0

CO

<Si CJ X!oH

(1)

+->CJM o

oj <u o > > o

in

vO o o

CO CM ^ fM ^ vO o O r-^ .—I

CO CO CO

CO

0) CO

.!H

—< U (U

(U (U

CO oO

M

O

O

O

O

bO do

fM (NJ

CO fM r- o o

in§ 00 ^ <

COo

0)o u

u o o o o

On

in Tt< fM vO

(M (M

00 in

CO CO

fMI I I—t r—1 < < in vD I I

COo

i3

a,

CJo

q;

CO

CO W
,500"

9,37"R

2.81"

FIGURE 1 - Tensile Specimen (1 in. = 2.54 cm)
95
Pull Rod

Stand Offs

0 2 4

1 I I L_J

SCALE, inches

0 5 10

J \ I SCALE, cm

FRP Insulators

Grip

ASTM

Clip Gage^

Bridge

CT Fracture Specimen

FRP Helium Dewar, vacuum insulated

FRP Nitrogen Dewar,

foam insulated

FIGURE 3 - Liquid helium fatigue cryostat

96
Removeable

Knife Edges

P

B = 1.50 in

W = 3.00 in

(1 in = 2.54 cm)

FIGURE 4 - Compact tensile specimen used for

fracture tests of Ti-5Al-2.5Sn

97
98
280x10

240

200

160

120

80

40

LONGITUDINAL, WARREN & REED (1963) • TRANSVERSE, THIS REPORT

TENSILE

YIELD

± ±

60 120 180

TEMPERATURE, K

240

150x10

CNJ

100

I/O LU

- 50

300 0

FIGURE 6 - Yield and tensile strength of Ti-5Al-2.5Sn

as a function of temperature.

9^
280x10*

240 -

200 -

160

120

80 -

40

T 1 INCONEL 750

TENSILE

YIELD

LESSMANN, et al

(1974)

• THIS REPORT

Sol Treat 1800° F,

Aged 1350°F. 8 h

Then Aged 1350Ho

1150°F. 10 h

Sol Treat 1800°

Aged 1625°F, 24 h

Then Aged 1300° F, 20 h

±

100 200

TEMPERATURE, K

150x10'

CSJ

100

UJ

a:I50

-J 0

FIGURE 7 Yield and tensile strength of Inconel 750

as a function of temperature.

100
.1 1 1 1 1 1 Jn

'O 60 120 180 240 300

^

TEMPERATURE. K

FIGURE 8 - Yield and tensile strengths of AISI 310 as a function of temperature.

101
280x10^

240

•r- 160

o.

CO

CO

UJ

on

\-

Ti-5Al-2.5 Sn

298 K

76 K

298 K LONGITUDINAL, WARREN & REED (1963)

O TRANSVERSE, THIS REPORT

X SAMPLE FRACTURED

- 150x10'

CM

100

CO

50

0.08 0.12

STRAIN

0.16 0.20

°

FIGURE 9 - Stress-strain curves for Ti-5Al-2.5Sn

102
280x10*

40

A-286

J.

Long . , To Fracture , Warren & Reed (1963)

Q Trans., Not Fractured,

This Report

0.08 0.16 0.24 0.32

STRAIN

FIGURE 10 - Stress-strain curves for A-286

ISOxlO'

CM

100

a:

- 50

0.40

Ho

103
STRAIN

FIGURE 11 - Stress-strain curves for Inconel 750.

104
210 T T

180-

30

AISI 310

X SAMPLE FRACTURED

O SAMPLE DID NOT FRACTURE

298 K

4 K

120x10'

90

CSI

60

30

0.15 0.30 0.45 0.60 0.75

STRAIN

FIGURE 12 - Stress-strain curves for AISI 310

105
STRESS INTENSITY FACTOR RANGE, AK (N/m^'/iT)

STRESS INTENSITY FACTOR RANGE, AK (ksi-in^^")

FIGURF ]3 - Crack growth data for T i - 5 A 1 - 2 . 5 S n

I Of.
STRESS INTENSITY FACTOR RANGE, AK (N/m^-m^)

5x10

-4

5 10

1

\—I ! I I I

10^

^ 10"^

o

>>u

<a.o

Q.

g 10-^

LiJ ZDC5

10

I I I I I I

I

A-286

1 1—I I It 1 I

I A -I

9:

• *

•* f

: I-

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A 76 K - 3 Specimens

4 K - 3 Specimens

I I I I I I I

10

_

- — 10"

_

10

- 5

2 5 10 5

STRESS INTENSITY FACTOR RANGE, AK (ksi-in^^)

FIGURE 14 - Crack growth data for A-286.

Jj

100

0)u

o

-a

CD<t

Q.o

Q.

ZD

107
10

-4

STRESS liTEMSITY FACTOR MMl, AK (N/Bp./iT)

10'

T

t I f ! i n I

I ?

' ; » f M 1 1

I I I It I I M I T

1 II I!

S

10-*

MCOKl 750

4 K

- 10

-4

m

A

§

A

10

-5

10

-6

10

-7 J I L

10

J I I i I I 50 100

STRESS INTENSITY FACTOR RANGE, AK (ksl-in^^^)

FIGURE 15 - Crack growth data for Inconel 750.

108 X
10

10-5 1-

10

10

-7

STRESS IHTEWSITY FACTOR RAM6E. AK (N/m./F)

10^

1 1 1 ? M M I "T 1—I I M M I

AISI 310

• 298 K A 76 K m 4 K

J I I I I I 1

1

10

10

8

1 1 1 I I I I I I n—I—I I I I I u

A

- 10

-4

— 10

-5

10

-6

50 100

STRESS INTENSITY FACTOR RANGE, AK (ksiin^/^)

FIGURE 16 - Crack arowth data for AISI 310.,

109
STRESS INTENSITY FACTOR RANGE, AK (N/mVT)

2 5 10®

STRESS INTENSITY FACTOR RANGE, AK (ksi in^^"^)

FIGURE 17 - Crack growth data for AISI 304,310,

and 316 stainless steels at 4K.

110
SPECIMEN EDGE DISPLACEMENT

FIGURE 18 - Fracture test records for Ti - 5A1 - 2 . 5Sn

Ill
1 PYLE et al

2 CARMAN et al

3 NBS data

- 90x10

Ti-5Al-2.5 Sn

NBS data for Ti-6A1-4V alloy

100 200

TEMPERATURE, K

- 80

- 70

60

o

t—

(

cr

50

40

a:

t30

300

FIGURE 19 - Temperature dependence of fracture toughness of

titanium alloys.

112
DISPLACEMENT, cm

0 0.02 0.04 0.06 0.08

DISPLACEMENT, in

FIGURE 20 - Fracture test records for A-286.

113
o

1

1 lo

1^ UD Lf5 CO ^UL/qi-UL *-|Vb931NI-r

o

X

114
FIGURE 22 -

100 200

TEMPERATURE, K

Temperature dependence of fracture toughness

for A-286.

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:106

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- 150

- 100

E

<N

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o

LlJ

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100

TEMPERATURE. K

FIGURE 25 - Temperature dependence of fracture

toughness for Inconel 750.

300

118
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NBSIR

SEMI-ANNUAL REPORT ON MATERIALS RESEARCH

IN SUPPORT OF SUPERCONDUCTING MACHINERY

ADVANCED COMPOSITES

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

Cryogenics Division

Institute for Basic Standards

National Bureau of Standards

Boulder, Colorado 80302

October 1974
Summary; Advanced Composites

The experimental work in this program will be conducted in two phases.

Phase I, already underway, will have the objectives a) to thoroughly check out

and refine experimental methodology, instrumentation and equipment and, b) to

obtain base data on commercial boron/epoxy and graphite /epoxy against which to

compare the results obtained in Phase II. Phase I work will be confined to

evaluation of uniaxial longitudinal and uniaxial transverse tensile strength

and moduli at 295 K, 77 K and 4 K.

Phase II will provide an in-depth characterization of the static mechanical

properties of uniaxial lamellae of boron/epoxy, boron/aluminum, graphite /epoxy

,

Kevlar 49/epoxy and glass/epoxy down to 4 K, emphasizing the latter temperature

as it is of most interest to superconducting technology. The program is designed

to yield data which may be used by designers faced with the problem of calculating

the properties of complex fiber layups in real engineering structures. Phase II

materials will consist of commercial boron/epoxy and boron/ aluminum, while the

glass, Kevlar 49 and graphite fibers will utilize an epoxy matrix optimized for

ryogenic use. Sufficient Phase II material has been ordered to permit thermal expansion

(contraction), thermal conductivity, dynamic elastic modulus characterization

v"it the materials.

A comprehensive review of the literature on the mechanical and thermal

properties of advanced-fiber reinforced structural composites at cryogenic

temperatures has been completed and is included in this report. This is Part II

of a two-part series. Part I reviewed the literature on glass-reinforced com- posites and was included in the ARPA Semi-Annual Technical Report May 1974.

The Part II review includes tensile, flexural and compressive strength and

moduli, interlaminar shear, ultimate tensile strain, bearing strength, dynamic

fatigue and impact, thermal expansion (contraction), thermal conductivity and

specific heat. Only uniaxial longitudinal and uniaxial transverse data are considered. Composite classes include boron/epoxy, boron/aluminum, graphite /epoxy

and Kevlar 49/epoxy. Data are presented for several hybrids (combinations of

reinforcing materials) for comparative purposes. The amount of data available

varied greatly among the composite types. In a few cases, a very large amount

of data were available for a specific property of a specific composite type (e.g.,

flexural data on HT-S/X-904 graphi te/ epoxy ) ; however, in most cases, available

data reflected the results of only a few tests.

124
The literature Indicates that the advanced- fiber composites have potential

application in cryogenic structures where a high modulus is required along

with excellent strength and (except for aluminum-matrix composites) low thermal

conductivity. The commercial boron/epoxy and boron/aluminum composites appear

to perform exceptionally well without optimization for cryogenic service. Results with graphite- reinforcement are less clear. Available data suggest some erratic behavior at cryogenic temperatures for these latter materials; however,

some successful practical applications of graphite /epoxy composites at cryogenic temperatures suggests the possibility that the erratic experimental data

may reflect problems of test method rather than inherent material deficiency.

Additional study is needed in this area. The relatively new Kevlar 49 fiber in epoxy matrices appears promising

for cryogenic work, although available data are minimal. The advantage of

this fiber is relatively low cost combined with a substantially higher modulus

than is obtained with glass reinforcement.

For most properties and most materials, available data do not extend below

77 K. No direct data (as distinct from that generated in pressure vessel evaluation

tests) exists for fatigue performance of advanced composites at cryogenic temperatures. Although there is no a priori reason to believe that the generally excellent

fatigue properties of such materials at room temperature will be degraded by

cooling, this parameter warrants further investigation.

Available data on polymeric-matrix composites are almost entirely confined to

epoxy matrices, although a small amount of data were available for a polyimide

matrix. The epoxies were for the most part conventional formulations not optimized for cryogenic use. While this does not appear to be a factor in the boron/

epoxy types, data suggest that the graphite and Kevlar 49 fiber composites may

benefit from matrix optimization.

125
Contents : Advanced Composites

Page

A. COMPOSITE TESTING PROGRAM

1. Introduction 127

2. Phase I; Preliminary Work . 128

2.1 Procedural Checkout 128

2.2 Acquisition of Initial Data 131

3. Phase II: Planned Testing Program 135

3.1 Materials Selection 135

3.2 Mechanical Test Matrix 138

3.3 Provisions for Additional Characterization 140

A. Future Work Requirements 140

B. MECHANICAL AND THERMAL PROPERTIES OF FILAMENTARY-REINFORCED

STRUCTURAL COMPOSITES AT CRYOGENIC TEMPERATURES -I I

Advanced Composites 143

126
COMPOSITE TESTING PROGRAM

1. Introduction

The composite testing program was finalized following completion of the

literature review, Part II of which preceeds this Section. This review led

to the following conclusions relative to the test program:

1) Much of the literature data are of questionable value due to the lack

of standard test procedures, inadequate numbers of specimen replications or to

poor experimental technique. For the most part, cryogenic data were incidental

to the main testing program with little attempt to optimize test procedures for

low temperature conditions

2) Much of the literature data are of less value than desired due to failure

to report the full characterization of the material under test, i.e., void fraction,

fiber density, processing details, location and type of fracture, etc.

3) Existing data, while useful in illustrating trends, are of relatively

little use to designers, as the data are insufficiently complete to permit prediction of crossply properties.

4) Existing data extend only to 20 K for glass, while most of the advanced

composite data terminate at 77 K.

5) The commercial boron/epoxies and boron/aluminum (6061) composites

appear to perform sufficiently well at cryogenic temperatures as to preclude

the need for optimization. However, the glass, graphite and Kevlar 49 composites

appear to benefit from the use of an epoxy matrix optimized for cryogenic use.

Based on these considerations, a two-phase testing program has been

initiated.

127
2. Phase I; Preliminary Work

This phase has been designed In particular response to Item (1) above.

Its Immediate objectives are to firmly establish the validity of the method

for tensile testing down to 4 K. At the same time, this phase will produce

mechanical property data on commercial boron/epoxy and commercial graphite/

epoxy which will be useful as a base line for comparison with results obtained

in the subsequent, more comprehensive program.

The composite materials used in this preliminary study have been donated

by two producers of commercial prepreg materials, all material being unidirectional:

Material Supplier Plies Orientation

'3.6 mill Boron/Epoxy

5.6 mil Boron/Epoxy

5.6 mil Boron/Epoxy

Craphi te/Epoxy

Graphite/ Epoxy

A

A

B

A

A

4

11

6

6

16

longitudinal

transverse

longitudinal

longitudinal

transverse

2.1 Procedural Checkout

The procedure for tensile testing was based on ASTM D 3039-71T, Tentative

Method of Test for Tensile Properties of Oriented Fiber Composites, with the

gripping system modified for cryogenic testing. All specimens are 28 cm (11 in)

long with a 15.2 cm (6 in) gage length. Longitudinal specimens are 1.27 cm

(1/2 in) wide, while transverse specimens are 2.54 cm (1 in) wide.

The prior experience of our laboratory and the available literature concur

in selecting strain gages over extensometers for precise strain measurement;

furthermore, determination of the overall Poisson's ratios requires transverse

strain data for which purpose strain gages must be used.

128
One of the major problems in mechanical testing of composites is

obtaining valid tensile fractures in the uniaxial longitudinal test mode,

i.e., obtaining fractures within the gage length rather than within the grips.

The problem is basically one of transferring the tensile load from the grips

into the filaments of the composite. Our approach to solving this problem

differs from ASTM D 3039 in that we do not plan to use reinforcing tabs.

Rather, we have chosen to use the grip system illustrated in Figure 1. In

this arrangement, a piece of 100 mesh stainless steel screen of 0.01 cm

(0.0045 in) wire is wrapped around each end of the specimen for a distance

of about 6.4 cm (2.5 in). The specimen ends are then sandwiched between two

sheets of 0.16 cm (1/16 in) Ti-6A1-4V titanium alloy which has been crossserrated with 10 teeth per inch on the sides contacting the specimen. Finally,

two plates of 0.48 cm (3/16 in) thick stainless steel pressure plates are

bolted onto the assembly. In this arrangement, the tensile load is transferred

only through the titanium part of the grip.

The pressure is applied by a row of six 8-32 machine screws on each side

of the grip. The torque on the screws is decreased systematically going from

the grip ends toward the gage section so as to introduce the load into the

specimen in a systematic manner. For example, with boron/epoxy material, a

satisfactory arrangement has been found to be 0.25 N • m (35 in-oz) of torque

applied to the four screws closest to the specimen end, 0.23 N • m (33 in-oz)

for the screws fifth from the end and 0.22 N • m (31 in-oz) for the screws

closest to the gage length. The grips and specimens are mated in an alignment

fixture prior to insertion into the tensile cryostat.

As an initial step in the procedural checkout, the alignment of the

specimen as mounted in the cryostat was determined by the multiple-gage techniqi

129
SERRATED

SURFACE

Figure A-1 Grip Assembly for Tensile Testing of Uniaxial

Composites at Cryogenic Temperatures

130
specified in ASTM D 3039-71T and was found to be in compliance with the stated

criteria. At the same time, the two separate strain gage balance, amplification

and recording systems as required for Poisson's ratio determination were checked

out and proven satisfactory.

2.2 Acquisition of Initial Data

The initial run was made using a 4-ply 5.6 mil boron/epoxy longitudinal

specimen from supplier A. This specimen was cycled through a portion of its

elastic region three times at room temperature (295 K) and the longitudinal

and transverse modulus was determined. The procedure was then repeated at

liquid nitrogen temperature (77 K) and in liquid helium (4 K) . The specimen

was stressed to failure on the fourth cycle at 4 K. The failure was a valid

fracture, occurring near the center of the gage length. The fracture is

illustrated in Figure 2.

The data obtained from these tests are illustrated schematically on

Figure 3 and is summarized on Table 1. The latter table also contains initial

ultimate tensile data obtained at 77 K on another specimen. This work confirmed

adequate sensitivity of the measurement system. As is usual for composite specimens

the initial cycle at room temperature produced a distinct hysteresis loop which

did not reappear during subsequent cycling. This initial hysteresis is conventionally attributed to isolated failure of some of the filaments, perhaps at

filament kinks.

The small 3 a values associated with the repeated determinations on

Table 3 confirm excellent repeatability of the test system. Unfortunately, a

lead to the transverse gage became disconnected during cooldown to 4 K,

preventing calculation of the Poisson's ratio for this temperature.

131
Figure A-2 Typical Tensile Fracture in 4-Ply Uniaxial

Longitudinal 4.6 mil Boron/Epoxy at 4 K. M = 3

132
LONGITUDINAL

GAGE

TRANSVERSE

GAGE

Figure A-3 Longitudinal and Transverse Stress-Strain Data Obtained

From a 4-Ply 5.6 mil Boron/Epoxy Specimen Tested in the Uniaxial

Longitudinal Mode at 298 K, 76 K and 4 K

133
The 2 73 KSI ultimate tensile strength measured at 4 K and the 256 KSI

at 77 K in the present work is somewhat higher than the average literature

values presented in the review section of this report. However, the present

experimental results are reasonably in line with the highest literature values

since, as noted in the review text, a 77 K value of 226 KSI was reported for the

(13) ultimate tensile strength of 4.0 mil boron/epoxy by Hertz, et al. . A similar

comparison of the temperature dependence of the longitudinal modulus values of

Table 1 with the literature averages presented in the review show the present

experimental data to be consistently about 2 x

10''"^ N • m ^ (3 x 10^ psi) higher

than previously reported.

The helium consumption of about 15 liters was rather high during this

initial test. It is expected that system improvements and operating experience

will significantly reduce the quantity required per test.

TABLE 1

INITIAL CRYOGENIC MECHANICAL PROPERTY DATA

ON 4-PLY 5.6 MIL UNIAXIAL LONGITUDINAL BORON/EPOXY

Temperature

Property 295 K 77 K 4 K

,ti '1 E^"" 10^ N • m~^ 238(1)* 249(1) .253(2)

10^ psi 34.5(0.2) 36.2(0.2) 36.6(0=3)

Y — 0.236(0.003) 0.256(0.030)

o^" 10^ N • m~^ — 17.7 18.8

10-^ psi — 256 273

E^"" 10"-^ - 7.8 7.4

*() = 3a values for successive modulus determinations on same specimen.

134
3. Phase II: Planned Testing Program

This Section describes the rationale for selecting the materials and

describes the specific composites to be included in the program. The rationale

for the specific testing program is described and the testing matrix is illustrated.

Finally, additional desired follow-up work is described.

3.1 Materials Selection

Composites having potential use in superconducting machinery are boronreinforced aluminum and boron-, graphite- glass- and Kevlar 49-reinforced

epoxies. The characteristics of each of these composite classes are summarized

on Table 2.

The literature reivew indicated that composites fabricated directly from

commercial boron/aluminum 6061 alloy "green" tape and from boron/epoxy "prepreg"

tape will have excellent strength and will show little temperature sensitivity on

cooling to cryogenic temperatures. These materials have therefore been ordered

as state-of-the-art commercial composites.

As fabrication of boron/aluminum composites requires specialized equipment,

this material has been ordered fabricated directly to the desired shape from a

primary supplier. The remaining composites have been ordered from a commercial

job shop specializing in supplying polymeric-matrix structural composite components

to industry. The latter materials will therefore be representative of commercial

production. The boron/epoxy composites will be fabricated from commercial 3-inch

wide prepreg tape. However, the glass, graphite and Kevlar 49 composites will

be produced with a NASA Resin 2 matrix, as the latter resin has been specially

developed to provide increased flexibility at cryogenic temperatures.

The selected composite test materials are listed in Table 3. S-glass

has been selected as it is the most common type used in high strength glass

135
TABLE 2

CHARACTERISTICS OF SELECTED COMPOSITE CLASSES

Composite Class

Boron/Aluminum

Boron/Epoxy

Graphlte/Epoxy

Kevlar 49/Epoxy

Glass/Epoxy

Characteristics

High strength, high modulus, limited

formabillty of finished sheet, high

thermal conductivity, high thermal

contraction, low electrical resistivity,

high cost

High strength, high modulus, low thermal

conductivity, low thermal contraction,

high electrical resistivity, high cost

Moderate strength, medium modulus, medium

thermal conductivity, very low thermal

expansion, moderate electrical conductivity,

moderate cost

Moderate strength, medium modulus, low

thermal conductivity, high thermal ex- pansion, high electrical resistivity,

low cost

Very high strength, very low modulus,

low thermal conductivity, low thermal

contraction, high electrical resistivity,

low cost

136
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composites. Type A (or AS) graphite fiber will be used as these are

relatively low cost fibers of intermediate strength and modulus and are

well characterized. Boron 5.6 mil diameter has been selected as this diameter

fiber is currently replacing the previously used 4 mil size.

3.2 Mechanical Test Matrix

The principle underlying the mechanical testing program is to provide

a basic static mechanical property characterization of a uniaxial lamella of

each composite type at 4 K. Characterization will also be undertaken at 77 K

and 295 K where such data are not already available. Highest priority will

be placed on the 4 K characterization.

From the stress-strain curves in longitudinal tension, transverse tension,

longitudinal compression, transverse compression and in-plane shear, the elastic

constants E^^, ^22' ^12 ^12 ^^^-^ determined along with the proportional

limits and ultimate strengths. These data may then be used by designers to make

a limiting static strength analysis of off-axis laminates made from the composite

of each basic lamella. This approach is based upon modern composite strength

theory and is currently extensively used in the aircraft and other industries.

The test matrix of Table 4 summarizes the test series required to obtain

this objective. A minimum of five valid test results must be averaged for

each parameter variation in order to establish the required confidence levels.

Priority will be given to obtaining longitudinal and transverse tensile and

in-plane shear data at 4 K, as these data by themselves suffice for design

approximations if it is assumed that the compressive strengths are equal to

those in tension. As the compressive strengths are in fact superior to those

in tension, this assumption leads to a very conservative limiting static

138
TABLE 4

MECHANICAL CHARACTERIZATION TEST MATRIX

Legend :

BFRP - boron fiber reinforced plastic TT - transverse tension

BFRM - boron fiber reinforced metal LT - longitudinal tension

GFRP - glass fiber reinforced plastic TC - transverse compression

CFRP - carbon fiber reinforced plastic LC - longitudinal compression

K49RP - Kevlar 49 reinforced plastic S - in-plane (intralaminar) shear

Composite Test Temperature and Number of Specimens

Type Tvoe 295 K 77 K 4 K

CI") BFRP^ LT 5 5 5

It TT 5 5 5

It LC 5 5 5

II TC 5 5 5

II BFRM^^^ LT

5

5

5

5

5

5

II TT 5 5 5

II LC 5 5 5

II TC 5 5 5

11 (3) CFRP^^

S LT

5

5

5

5

5

5

11 TT 5 5 5

It LC 5 5 5

It TC 5 5 5

II S 5 5 5

K49RP LT 5 5 5

II TT 5 5 5

II LC 5 5 5

It TC 5 5 5

It GFRP^^^

S LT

5

5

5

5

5

5

II TT 5 5 5

It LC 5 5 5

II TC 5 5 5

II S 5 5 5

(1) 5.6 mil commercial boron/ epoxy

(2) 5.6 mil commercial boron/ 6061 aluminum

(3) Type A (or AS) graphite/NASA Resin 2

(4) Kevlar 49 /NASA Resin 2

(5) S-901 glass /NASA Resin 2

139
strength prediction and probable overdesign of components. Consequently,

the next priority will be given to obtaining valid compressive strength data

at 4 K. Following this work, the static property data at 77 K and 295 K will

be obtained, the latter data being useful in designing components exposed to a

range of temperature.

In addition to the defined 4 K characterization studies, a few ad hoc

tests will be performed to obtain preliminary data on the extent to which the

strength and modulus of these composites may become degraded by repeated

thermal cycling (thermal fatigue) over the 295 K-4 K temperature range.

3.3 Provisions for Additional Characterization

Sufficient material is being obtained in each composite type to permit

thermal characterization of the materials as needed to verify existing data

and to extend existing data to 4 K. Such characterization includes dimensional

changes on cooling and the temperature dependence of thermal conductivity.

Additionally, sufficient material is being obtained to facilitate a

study of the feasibility of using dynamic methods of determining the significant

engineering elastic constants in lieu of the conventional mechanical method.

Direct comparison of values obtained by each method on the same composite will

be possible. The dynamic modulus will be investigated by both the pulse-echo

and the resonance methods

4. Future Work Requirements

Aside from the ad hoc thermal fatigue tests mentioned above, the present

program is confined to characterizing static mechanical properties. Performance

of composites under dynamic loading conditions at cryogenic temperatures also

must be determined before such materials can be used with confidence in rotating

components of superconducting machinery. Unfortunately, djmamic fatigue testing

140
at cryogenic temperatures is exceedingly expensive. Equipment must be tied

up for a long time, the cryogen consumption is high and the anisotropic nature

of composites requires analysis of several fiber orientations. Cryogenic

fatigue testing must therefore be confined to those composites appearing most

applicable for specific superconducting machinery components and such tests

must be carefully planned to produce the maximum amount of useful data.

141
142
MECHANICAL AND THERMAL PROPERTIES OF FILAMENTARYREINFORCED STRUCTURAL COMPOSITES AT CRYOGENIC

TEMPERATURES-II: Advanced Composites

Maurice B. Kasen

Cryogenics Division, NBS-Institute

for Basic Standards

Moulder, Colorado 80302

ABSTRACT

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

and thermal properties of advanced-fiber reinforced structural composites at

cryogenic temperatures. The objective is to provide an understanding of the

general magnitude of the property values obtainable within the cryogenic temperature range, to provide a feel for the general ranking of various specific

composite types with regard to specific properties, and to impart an understanding of the temperature dependence of the property of interest . A comprehensive Bibliography and Bibliography-Property Cross-Reference is included.

This is Part II of a two-part series. Part I considered glass-reinforced

composites.

This research was supported by the Advanced Research

Projects Agency of the Department of Defense under

ARPA Order No. 2569.
SYMBOL NOMENCLATURE

The following symbol nomenclature is used in this report:

a^^ - tensile ultimate strength

ty a ^ - tensile yield strength

E^ - initial tensile modulus

E2 - secondary tensile modulus

tu

e - tensile ultimate strain

n^^ - tensile fatigue failure stress

o^^ - flexural ultimate strength

E^ - initial flexural modulus

E2 - secondary flexural modulus

cu

a - compressive ultimate strength

E - compressive modulus

si a - interlaminar shear strength

by

a - bearing yield strength

a^^ - bearing ultimate strength

- impact strength

A, - thermal conductivity

AL/L - thermal contraction

- specific heat

p - density

psi - pounds per square inch

KSI - psi X 10^

2 N • m - Newtons per meter squared (Pascal)

J • K '"K - Joules per kilogram - Kelvin

g

W • m~'^K ^ - Watts per meter-Kelvin

144
INTRODUCTION

The objectives of this review are fourfold: a) to provide the designer

with a feel for the general magnitude of property values that may reasonably

be expected from a given category and class* of composites within the cryogenic

range, b) to provide him with a feel for the ranking of specific composite

classes with regard to a specific property, c) to impart a feel for whether

the property of interest is likely to increase, remain unaffected, or decrease

with lowering of temperature, and d) to define those areas in which additional

data are needed and to define the direction that future work should take in this

area. Readers with more specific interests are referred to the Bibliography and

Bibliography-Property Cross-Reference for retrieval of specific documents. The

Bibliography in this report is similar to that in Part I and includes references

to both the glass- and advanced-fiber reinforced composite literature. The scop

of the literature survey has already been described in Part I.

Several differences will be noted between the presentation of the data

in this work and in Part I. First, because of the wide variation in properties

among the advanced fibers (and, therefore, among the composites in which they

are used), separate data are presented on each specific reinforcement type.

Secondly, as no data were available on woven-cloth advanced-fiber composites

at cryogenic temperatures, present data are restricted to uniaxial-transverse

lajmps. Finally, the overwhelming majority of the data for advanced composite

polymeric systems are reported for epoxy matrices, in contrast to the variety

of matrix types reported for glass-reinforced systems.

* We define a composite category by the general reinforcement type, e.g., glass

fiber or advanced fiber (graphite, boron, etc.). We subdivide the category into

composite classes by the general matrix type, e.g., glass/polyester or graphite/

epoxy. We further subdivide the class by referring to a composite type when a specific reinforcement/matrix combination is specified, e.g., HT-S/X-904 epoxy.

145
In other respects, the data presentation follows that used in Part I, i.e.,

literature property values are presented in graphical form as a function of temperatures at 295 K, 200 K, 77 K, 20 K and 4 K. The absence of a data point for

a given temperature indicates failure to find significant data. References are

given for each plotted curve.

The reader is cautioned that curves presenting data averaged from several

sources may have a considerable scatter band associated with them. We discuss

the range of values associated with such curves and emphasize those specific types

for which the best values were reported.

The author has attempted to make this review as comprehensive as possible.

Nevertheless, the complexity of the subject makes it unavoidable that some data

worthy of inclusion have been inadvertently overlooked. The author would appreciate

having such omissions brought to his attention. The author also wishes to emphasize that the data presented in the review reflect the published results of the

cited authors. These data have not been experimentally verified by NBS, and the

conclusions and evaluations presented herein do not imply approval, endorsement,

or recommendation of any commercial product by NBS.

In considering the mechanical property data, the reader should be aware that

there exists no universally accepted method of determining most of these properties.

It is often difficult to obtain valid uniaxial longitudinal tensile fractures

because of the difficulty of transferring the load from flat specimen grips into

the fibers of the specimen. Failure within or adjacent to the grips is not uncommon. This problem is particularly acute when testing CFRP materials due to

the small fiber diameter. Problems also arise in uniaxial longitudinal compression testing where an unsupported specimen may fail by "brooming" of its ends

or by column buckling. However, such premature failures result in property

values lower than the true values and tend to bias the data accordingly.

146
The reader should also be aware that the discussion of properties of the

composites included in this review does not take into consideration the effect

of variations in fiber/resin content 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 reflect actual values and trends reported for

specific composites. Controlled variations in many of the properties are obtainable in practice by specific variation of the fiber content of the composites.

Finally, the reader should be aware that composite technology is developing

so rapidly that some of the data presented in this review may not reflect the

current state of the art. New reinforcing fibers, inproved composite fabrication

techniques and refined test methods will very likely result in overall improvement in both the level and the consistency of composite mechanical properties.

For the reader unfamiliar with the development of advanced composite

technology, it will be useful to establish a perspective on the field. The two

primary reinforcement systems are boron filaments produced by vapor deposition

of boron on a very fine tungsten wire substrate and graphite fiber produced

by graphitization (pyrolysis) of an organic precursor fiber. Additionally,

a proprietary organic fiber called Kevlar 49* is currently receiving attention.

The boron filaments are normally produced to 4 mil diameter, although 5.6 mil

diameter filaments are also coming into use. Boron fiber coated with a thin

layer of silicon carbide is produced under the trade name of Borsic* and is

reported to have improved interfacial bonding to certain matrix types. The

graphite fiber field is much more complex, as it is not only possible to produce

* The use in this paper of trade names of specific products is essential to the

proper understanding of the work presented. Their use in no way implies approval,

endorsement or recommendation by NBS

.
such fibers from different precursor materials,* but it is also possible to

vary the production process to produce fibers differing greatly in modulus

and strength. As the properties obtainable with graphite reinforcement may

approach those with boron reinforcement at a somewhat lower cost, development

of graphite fibers has proceeded at a very rapid rate until today the user is

confronted with an abundance of fiber choices, many of which are not well

characterized and many of which will disappear to be replaced by newer types.

At the present time, the potential user of these materials would be wise to

restrict his interest to those types whose behavior is reasonably well known,

barring compelling reasons for doing otherwise.

Graphite fibers may be produced with elastic moduli varying from 25-75 x

10^ psi (17-50 X lO"*"^ Pa), with strengths varying inversely to moduli. In the

present report, these fibers are classified according to low, medium or high

modulus, the differentiation being <40, 40-60 and > 60 x 10 psi, respectively.

As composites are frequently used where weight is critical or where high

specific strengths are required, typical composite densities have been summarized

in Table 1.

STATIC MECHANICAL PROPERTIES

For most mechanical properties, the data for the CFRP composites has been

separated from those of the other advanced composites. This reflects the larger

amount of available data on CFRP materials and the large variety of graphite

fibers for which data have been published.

Composite Tensile Strength

Figure 1 summarizes the available data on the temperature dependence of

the ultimate tensile strength of graphite- reinforced epoxy composites from

* The two most common are rayon and polyacrylonitrile (PAN).

148
room temperature into the cryogenic temperature range for the uniaxial longitudinal

and uniaxial transverse orientations. A surprisingly large amount of data were

found in the literature for CFKP materials—58 separate reports of test data,

each report being the average of several tests on a given composite. Undoubtedly,

this large effort reflects the desire to exploit the relatively low cost of

graphite fibers. Unfortunately, the available data summarized on Figure 1

indicate that CFRP composites suffer significant strength losses upon cooling

to 77 K and, as was observed with the glass- re in forced materials in Part I, the

strength behavior below 77 K appears to become erratic.

Figure 1(a) indicates that the uniaxial tensile strength obtainable in CFRP

composites is about 30% of that obtainable with glass-fiber reinforcement in

the cryogenic range. The ultimate tensile strength of composites tested in

the uniaxial longitudinal mode should be fiber controlled; hence, it would be

expected that the tensile strengths would inversely follow the modulus of the

fiber. This trend is not reflected in the data of Figure 1(a) —indeed, the

averaged data indicate the converse, higher strengths are associated with

higher modulus fibers and lower strengths with lower modulus fibers. However,

a more detailed examination of the data averaged into these curves shows that

the expected correlation does exist if one considers only the highest values

reported for each modulus range. Thus, the highest overall ultimate strength

13

at 77 K was 147 KSI reported for the low-modulus HT-S fiber in X-904 resin,

the next highest was 130 KSI reported for HMG-50 in an ERLB polyblend]"^

g for Modmor I and for Samco 360 in modified ERL 2256, all medium modulus

fibers, while 126 KSI was the maximum reported for the high-modulus Thomel

75 fiber in ERLB 4617^''" This suggests that the higher strengths associated

with the lower modulus fibers may have been lost in the averaging process due

149
to variations in composite quality or perhaps due to difficulties in obtaining

valid tensile fractures. The latter is a prime possibility, as it is particularly

difficult to obtain valid tensile fractures in uniaxial CFRP specimens loaded in

the fiber direction. Additional evidence that the test method is affecting the

results is found in the work of Simon and Larsen"*"^ who report almost diametrically

opposite temperature dependence of the ultimate tensile strength in HT-S reinforced

NASA Resin 2* and in an HT-S reinforced 4617 polyblend when each was tested first

as flat tensile coupons and subsequently as NOL (Naval Ordnance Laboratory) ring

specimens.

The uniaxial longitudinal tensile strength data for other types of advanced

composites are presented in Figure 2(a). The data for PRD 49/epoxy, boron/epoxy,

boron (Borsic) /aluminum, ** stainless steel /aluminum and the hybrid Borsicsteel/aluminum and Borsic- titanium/aluminum composites all have higher absolute

values of tensile strength and retain their strength to lower temperatures than

do the CFRP composites. The hybrid Borsic-s teel/aluminum composite contained

4.2 mil Borsic fibers in the tensile direction of 6061 aluminum and stainless

13 wires in the transverse direction. The hybrid Borsic-titanium/alumlnum

composite contained 4.2 mil Borsic fibers in the tensile direction with p-3Ti

13 foil interleaved between the boron lamellae. The latter composites are therefore not strictly tiniaxial; however, they have been included to illustrate the

interesting possibilities of hybrids. Of the two remaining uniaxial composites,

the HT-S/polyimide (Skybond 703) material displayed a comparatively low strength

and showed a significant decrease of strength upon cooling. Conversely, a very

* An^-epoxy resin formulation optimized for cryogenic service by Soffer and

Molho and consisting of Epon 828/DSA/EMPOL 1040 /BDMA in proportions 100/115.9/

20/1 pbw.

** Data for Borisc and boron fiber composites have been combined except where

differences in reported values justified separating the data.

150
sharp rise in strength on cooling was reported for, the HT-S/epoxy-phenolic

(HT-424 Primer), but its overall strength was not high.

The stainless steel reinforced composite contained NS-355 stainless wires

in a 2024 aluminum alloy matrix. This type of composite is available commercially

on special order; however, it has not received wide acceptance in view of the

wider availability of boron reinforced aluminum which has similar strength

properties.

Boron /aluminum composites are available commercially with either 4 mil or

5.6 mil boron or Borsic reinforcement. The literature values on the uniaxial

tensile strength at 77 K ranged from 163-202 KSI, with highest values reported

13 for 5.6 mil Borsic/6061. Data were not available for any other aluminum alloy

matrix. Boron/epoxy composites have been developed to an even higher degree

and are commercially available in the form of prepreg tape, i.e., with the

plastic matrix partially cured to facilitate component fabrication and to

improve composite quality. The literature values for boron/epoxy composites

tested in the longitudinal direction at 77 K varied from 167-226 KSI, the highest

13 value being reported for the commercial SP-2 72 product.

PRD 49 (Kevlar 49) is a relatively new type of organic fiber. Present

data indicate that this fiber in an epoxy matrix is capable of developing tensile

strengths comparable to that developed with graphite fibers at cryogenic tem59

peratures. Hoggatt reports NOL ring ultimate strengths of 146 KSI for PRD 49-1

in a NASA Resin 2 matrix and 183 KSI in an ERLB 4617 matrix at 20 K. There

is some evidence of a slight decrease in strength on cooling to 77 K; however,

available data indicate that the ultimate tensile strength of this material is

relatively independent of temperature within the cryogenic range.

The hybrids Borsic-steel/aluminum and Borsic- titanium/aluminum complete

the group of advanced con^josites that have reasonably high strength in uniaxial

151
tension. The lower strength of these hybrids compared to the conventional

boron/aluminum composite reflects the reduced density of boron fibers in the

direction of loading.

'In uniaxial transverse tension, Figure 2(b) shows the hybrid Borsicti tanium/ aluminum composite to have higher strength than the conventional

boron/aluminum type, which is precisely the reason for creation of the hybrid.

In this case, the benefit of the interleaved titanium foil increases the transverse

tensile strength at 295 K and 200 K; however, the effect appears to diminish

rapidly as the material is cooled to 77 K. While the transverse strength of the

boron/epoxy composite is much lower than that of the metal-matrix materials;

values are still in excess of those developed with the graphite fiber reinforcement.

The literature reported a transverse ultimate strength of 14.2 KSI for 4.2

13 mil Borsic/6061 aluminum and a somewhat higher value of 24.9 KSI for 5.6 mil

13 boron reinforcement, all at 77 K. It is probable that this strength difference

is real; however, available data also indicate that the observed difference may

be due to residual stresses inherent in the small diameter fiber rather than

to the presence or absence of a silicon carbide coating. The reader is referred

to the section on ultimate tensile strain for further discussion of this subject.

Composite Tensile Modulus '

The primary reason for development of advanced composites is the high

modulus obtainable with the newer types of fiber reinforcement. In contrast

to a maximum of about one million for glass- re in forced composites in the fiber

direction. Figures 3(a) and 4(a) show that moduli ranging from 30-40 x 10^

psi are obtainable with several of the advanced fiber composites. Thus, while

glass- reinforced composites may at their best equal the modulus of aluminum.

152
the modulus of composites using advanced fibers may equal and often exceed

that of steel.

Looking first at the graphite /epoxy composites, Figure 3(a) shows that

the modulus of the composite does, on the average, reflect the modulus of the

fiber when tested in the uniaxial longitudinal mode. The available data indicate

that the modulus of the Thornel 75 composite (the only representative of the

high-modulus fibers) declines to about that of the medium modulus fibers at

77 K. With this exception, the averaged data indicate that the moduli in the

fiber direction slightly increases on cooling to 77 K. No correlation is

observed between fiber modulus and transverse composite moduli in Figure 3(b)

;

indeed, none is expected, as the latter is controlled by the properties of the

matrix and by the fiber-matrix interfacial bond strength.

A comparison of Figures 3(b) and 4(b) shows that the uniaxial transverse

tensile moduli of graphite /epoxy composites is much lower than that developed

by the other advanced composites over the entire cryogenic temperature range.

A closer look at the literature data from which Figure 3(a) was prepared

revealed that among the medium modulus fibers, the modulus at 77 K was a respectable 39-43 x 10 psi for Samco 360, the highest value being reported with

g a modified ERL 2256 resin. (Note that these same fibers developed excellent

uniaxial tensile strengths.) A distinctly lower range of 26-33 x 10 psi was

reported for Thornel 50, highest values being reported for a modified ERL 2256

g composite. Among the low-modulus fibers, HT-S (the most tested fiber) produced

6 13 15-10 x 10 psi, highest value being reported with X-904 resin, while

17-22 X 10^ psi was the reported range for HMG-25, the highest values being

g obtained in a NASA Resin 2 matrix.

153
As the transverse modulus properties are matrix dominated, Figure 3(b)

shows significant differences for HT-S fiber in different epoxies. The

data illustrated are for the same composites as appear in Figure 1(b). Here,

the lower moduli of the HT-S/NASA Resin 2 and the HT-S/4617 polyblend are due

to the additions of f lexibilizers and elastomers to the epoxy resin while the

highest modulus was developed with the conventional X-904 matrix. In general,

the transverse modulus increases with decreasing temperature. A lower modulus

value was reported for the high-modulus Thornel 75 fiber in ERLB 4617 than

for HT-S/X-904.

Leaving the graphite- reinforced composites and turning to those reinforced

with other advanced fibers, we observe on Figure 4(a) that the "uniaxial longitudinal moduli reported for the boron/epoxy, the boron/aluminum and the hybrid

Borsic-titanium/aluminum and Borsic- steel/ aluminum composites all cluster

around 30-35 x 10 psi, with little temperature dependence. The PRD 49/epoxy

modulus is reported to be significantly lower, at 16-18 x 10 psi, but undergoing a significant rise between 77 K and 20 K. In general, these uniaxiallongitudinal moduli values compare favorable with those of the graphite-epoxies.

The uniaxial transverse moduli are another matter. As seen in Figure 4(b)

the transverse moduli of the aluminum-matrix materials are far higher than

those of the graphite/epoxies and are, in fact, almost twice as high as that

of unreinforced 6061 aluminum. The transverse moduli of the boron/epoxy is

much lower than that for the same fiber in an aluminiim matrix but is still

almost twice that of the graphite/epoxy materials. These data reflect the

contribution made by fiber-matrix interface to the transverse modulus.

Looking in more detail at the uniaxial longitudinal data, we find that the

boron (Borsic) /aluminum data ranged from 30.4-36.2 x 10^ psi with little temperature dependence. These data combine values obtained from fibers of 4 mil

154
and 5.6 mil in the two production variants, as no significant difference

was reported for these materials. The boron/epoxy represents average

81 data reported by Nadler, et al. on SP-272 and Narmco 5505 (data range was

not available). The hybrids, Borsic-titanium/aliominum and Borsic-steel/

aluminum were the same composites discussed in reference to the tensile data

of Figure 2(a). The slightly lower uniaxial tensile modulus of the hybrids

compared to that of the conventional boron /aluminum reflects the lower density

of fibers in the stress direction.

The PRD 49 data reflect input from both an ERLB 4617 and a NASA Resin

2 matrix^^'^^ The ERLB 4617 composite produced a slightly increased tensile

modulus at all temperatures, its value being 18.5 x 10 psi at 77 K compared

to 17.5 X 10^ psi for the NASA Resin 2. The large rise in modulus between

77 K and 20 K was reported for both matrices.

Considering the transverse data of Figure 4(b) in more detail, we observe

that the three aluminum matrix composites have similar moduli, about 20 x 10 psi.

The boron/aluminum data was again the average of data obtained from 4.2 mil

Borsic and 5.6 mil boron, similar values being reported for each variant. The

boron/epoxy data of Figure 4(b) again reflect the average values obtained

81 from the commercial SP-272 and Narmco 5505 materials.

Composite Flexural Strength and Modulus

Less data were available on the temperature dependence of the flexural strength

and elastic modulus than for tensile properties--f urthermore , most of that

available was reported for graphite/epoxy composites. Figure 5 summarizes

the flexural strength data for the latter, while Figure 6 summarizes the

available data for other types of reinforcements or matrices. The minimal

available data on flexural modulus properties appear in Figure 7.

155
Comparison of Figure 5 with Figure 1 shows that the graphite/epoxy flexural

strength data spans a much greater range of values than does the tensile

strength in both the longitudinal and transverse directions. The expected

higher strength of the lower modulus fibers is more in evidence for the

flexural test than it was for the tensile test mode.

By far the largest amount of data were available for HT-S/epoxy

composites; furthermore, such composites had the highest reported longitudinal

flexural strength of all the CFRP materials for which data were available.

Strengths were reported over a 113-234 KSI range at 77 K. This. is significantly

lower than the 325-470 KSI range reported for uniaxial glass-epoxies in Part I.

These data included test series designed to investigate cure cycles, aging

12 13 58 effects and environmental effects. ' ' Most of the work was done with

E-350, X-915 or X-904 matrices for which the average flexural strengths

13

at 77 K were reported to be very similar, varying only from 167 to 178 KSI.

It is beyond the scope of the present paper to discuss in detail the effect

of composite processing procedures on the resultant mechanical properties.

Nevertheless, the available data on the HT-S fiber composites remind

us that processing variables are important; in particular, those variables

that affect the void content. Detailed examination of the data shows that in

84 of the 89 reported test series, the flexural strength of composites made

with HT-S fiber declined markedly on cooling to 77 K. The five series

for which strength increases were noted were reported by Maximovich,

12

et al. to have occurred concomitant with a change to a vacuum-bagging

method of fabrication. This suggests that at least a part of the observed

drop in flexural strength on cooling may be due to the presence of voids

in the composites.

156
Continuing on with the other fibers, we observe on Figure 5(a) that

the order of decreasing strength is Modmor II, GY-70, and Fibralloy 300,

arriving finally at a group conprising the lowest strengths and consisting

of Courtaulds HM, HMG-25 and Thomel 50 fibers. The flexure strength

of this latter group is only about 1/3 that of the HT-S fiber composites.

Among these other fibers, most data were available for the high-modulus

GY-70, for which the spread of values at 77 K was 85-133 KSI, the highest

13 being reported in X-904 epoxy resin.

The clear separation between the transverse flexural strengths of the

low-modulus HT-S composites and the high-modulus GY-70 composites in

Figure 5(b) suggests that there are differences in the fiber -matrix bond

strength between these two fiber types, the high-modulus fiber having the

poorer bonding. Note that X-904 and 1004 epoxies were used with both fiber

types

Before leaving the data of Figure 5, it is of interest to consider whetlicr

or not the literature data supports the contention that the relatively inexpensive

flexure test may be used to obtain comparative ranking of composite strength

in lieu of the more expensive tensile test. A visual comparison of the data

on Figure 5(a) with that of Figure 1(a) is hardly convincing; however,

such a comparison is nebulous because these graphs present data from different

composite types. To clarify this question, the literature data were examined

more closely and a comparison of ranking of strength in longitudinal tension

versus longitudinal flexure was made at 77 K for eight specific composite

types (same fiber and matrix) for which data were available. The result is

summarized in Table 2, and suggests that the flexure test results are indeed

not a very good measure of the relative tensile strengths of graphite /epoxy

composites at 77 K. Insufficient data existed for comparisons with other

157
mechanical properties; however, the good performance of the HT-S fiber

composites in inter laminar shear (Figure 11) suggests that the flexural test

may correlate more closely with this parameter.

Flexure data on other types of advanced composites were somewhat

meager. Figure 6 does indicate that the uniaxial longitudinal flexural

strength of boron/epoxy in the form of the commercial SP-272 product is

much higher than that of the graphite/ epoxies, still lower than that developed

with glass reinforcement. A quite rapidly declining strength is observed

on cooling to 77 K. The latter effect was not observed in the comparable tensile

data (cf. Figure 2(a)). Data were available for two polyimide composites,

one with the high-modulus GY-70 fiber and one with the low-modulus HT-S fiber.

Of these, the HT-S developed the higher strength. A small decrease in flexure

strength is evidenced in these composites 77 K, but the magnitude is less than

that reported for the uniaxial tensile strength. Finally, the HT-S/epoxy-phenolic

(HT-424 Primer) appears to posses reasonable flexure strength with little

temperature dependence in contrast to the comparatively low strength and

significant temperature dependence reported for the same composite in longitudinal tension (cf. Figure 2(a)).

Very little data were available on the flexural modulus. Data for the

longitudinal mode appearing on Figure 7 reflect the modulus of the reinforcing

fibers and appear to agree reasonably well with the moduli of the same composites

in tension. The Thomel 50 data is an average of data from two matrix types,

g a modified ERL 2256 epoxy and for NASA Resin 2. A slightly higher modulus

was reported for the latter at both 295 K and 77 K.

Coiq>osite Compressive Strength and Modulus

Data on the compressive properties of advanced composites at cryogenic

temperatures were also relatively sparse. Those which were available appear in

158
Figures 8 and 9, combining the graphite/epoxy data with those of the other

composite types.

These data show that the advanced composites have excellent compressive

strength properties. In contrast to the tensile and flexural strength, the

compressive strengths and moduli of the advanced composites are significantly

higher than those of the glass-reinforced composites. Comparison of Figure 8

with Figures 1 and 2 also shows that the compressive strengths are 100-200%

higher than the tensile strengths for the same composites.

The highest uniaxial longitudinal compressive strength was reported

81 for the commercial boron/epoxy types SP-272 and Narmco 5505. The very

large strength increase on cooling to 77 K is most impressive for this composite

type. The boron/aluminum has a lower, but still respectable, uniaxial compressive

strength. The latter data reflect input from both 4.2 mil Borsic and 5.6 mil

13 boron fiber composites. A slightly higher strength was reported for the

5.6 mil composite, but not sufficiently to justify a separate plot. Again, the

Borsic-titanium/aluminum and Borsic-steel/aluminum hybrids developed

strengths somewhat below that of the conventional boron/ aluminum, reflecting

the decreased density of boron fibers in the stress direction. Lowest uniaxial

longitudinal compressive strengths were reported for the HT-S/X-904

graphite/epoxy composite; however, even this composite developed a compressive

strength about twice that developed in tension. Furthermore, the compressive

strength of the graphite/epoxy composite was reported to increase with cooling,

in contrast to the tensile behavior.

The available data on the transverse compressive strengths of boron/epoxy,

boron/aluminum and graphite/epoxy composites are compared on Figure 8(b),

showing these composites to be in the same relative order of strength as in

the longitudinal test mode. Again, the commercial SP-272 and Narmco 5505

159
boron/epoxy products indicate a substantial Increase in compressive strength

on cooling. Lesser increases are observed in the boron/aluminum and in the

HT-S/X-904 CFRP.

All else being equal, the compressive modulus of a composite should be

the same as the tensile modulus of the same composite. This is found to be

generally true for the HT-S/X-904 CFRP material, for the boron/epoxy for

the Borsic /aluminum composites and for the Borsic/st eel/aluminum hybrid, as

may be seen by comparing the compressive moduli values of Figures 9(a)

with the tensile moduli values of the same composites on Figures 3 and 4. The

Borsic-titanium/aluminum data are at variance with this principle, with some

10 psi higher modulus being reported in compression than in tension. The

boron/aluminum and the boron/aluminum hybrids again show the highest longitudinal moduli with the boron/ epoxy only slightly lower. The modulus of the

HT-S/X-904 CFRP composite is again much lower than that of the boron-reinforced

materials. There does not appear to be a significant temperature sensitivity

of this parameter. Also, as in the tensile case, the boron/aluminum demonstrates

a clearly higher transverse compressive modulus as compared to the boron/epoxy or

to the HT-S/X-904 CFRP composite.

Composite Interlaminar Shear Strength

Interlaminar shear strength is the resistance to failure of a composite

subjected to pure shear in the plane of the fibers, i.e., the shear strength

between two lamellae in a laminate along the plane of their interface.* Interlaminar shear is measured either by the guillotine method in which shear is

forced by the imposition of opposing but offset cuts in the sides of a flat

tensile specimen or by the short-beam method in which the specimen resembles

* Not to be confused with intralaminar shear which refers to shear within a single

lamina without reference to adjacent lamellae (sometimes referred to as in-plane

shear)

160
a flat flexural specimen but is designed to fail by shear on the central layers

of the composite. The guillotine method is the most simple and is reputed to

produce less scatter; however, the values obtained are usually higher than

those obtained with the short-beam method. The latter is considered to be the

more reliable method and is the most widely used. A modification of the shortbeam test is used with filament-wound NOL ring specimens; here, a short

section of the ring (concave downward) is substituted for the flat specimen.

Unfortunately, the results of the NOL segment and the flat short-beam tests are

not comparable.

The interlaminar shear properties are believed to strongly affect structural

integrity of a composite, particularly in compression loading and in flexure.

Like the flexural test, interlaminar shear tests simultaneously evaluate several

parameters of the conqjosite including resin strength, resin-fiber bond strength,

filament distribution and matrix void content. Such tests are therefore often

used to evaluate overall composite quality.

Because this test is a measure of composite quality, a wide range of

interlaminar shear strength values may be obtained with a single composite

t5?pe. An example of this is given in Figure 10(a) where a frequency histogram

depicts the range of longitudinal interlaminar shear strength reported in the

literature for the CFRP composite HT-S/X-904 at 77 K. A relatively substantial

amount of data are available for this composite, due in large part to a comprehensive

12 study of the effect of processing variables by Maximovich et al. Such

studies are of great value in perfecting fabrication methods; however, it is

evident from the strength distribution on Figure 10(a) that discussion of average

literature values of interlaminar shear strength (as has been done with other

properties discussed in this paper) would be relatively meaningless. More

useful would be knowledge of the maximum value obtainable in a "perfect"

161
composite of a specific type and the extent to which various defects degrade

the property. While available data within the cryogenic temperature range

do not permit such a detailed analysis, it is of interest to look at the higher

values reported for each reinforcement type in order to get a feel for what the

limiting interlaminar shear strength might be.

Before pursuing this objective, it is worthwhile to analyze the extensive

HT-S/X-904 data a bit further to determine what this relatively large body of

data can tell statistically about the temperature dependence of interlaminar

shear strength. From the frequency histogram of Figure 10(b) we observe that

the reported change in this parameter upon cooling from 295 K to 77 K approximates

a normal distribution around zero change, suggesting that interlaminar shear

strength is relatively independent of temperature for this composite over this

temperature range.

Figures 11 and 12 reflect the upper bounds of uniaxial longitudinal interlaminar shear strength obtainable with various types of advanced composites

based on available literature. In preparing these Figures, the highest five

reported values were averaged for the HT-S/X-904 data, while the data for

the other fibers reflect the average of all available data when the data spread

was minimal or, in a few cases, the selection of the higher of two groups

of data when such a grouping was obvious. Curves based on only one set

of data are identified by an asterisk.

Among the CFRP composites, the data suggest that the highest interlaminar

shear strengths will be obtained with the low-modulus HT-S fiber. These data

13 include three evaluations with X-904 resin by Hertz, one evaluation with

58 Ciba 8183, also by Hertz and one evaluation of NASA Resin 2 by Larsen,

et al!"^ The 77 K values ranged from about 17,500 psi for the X-904

composites to 21,000 psi for NASA Resin 2, indicating that such high values

162
are not unreasonable for good quality composites made with this fiber. The

room temperature data for the X-904 and Ciba 8183 specimens ranged from 13,000

to 16,000 psi; however, the NASA Resin 2 specimens developed only 7500

psi at this temperature, underlining that this latter resin sacrifices elevated

temperature properties for desired performance in the cryogenic range. Figure

11 also indicates a further increase in strength of HT-S composites upon cooling

to 20 K. The latter data lacks literature corroboration, relfecting only the

data published by Larsen, et al}^ with a NASA Resin 2 matrix.

Figure 11 indicates an appreciable increase in interlaminar shear strength

with lowering of temperature for the selected HT-S composites in contradiction

to the negligible dependence inferred statistically by Figure 10(b) for the

same type of composite. As there is no reason to question the validity of

either of these data, it would appear that processing variables may be affecting

the temperature dependence, high-quality composites increasing their interlaminar

shear strength at lower temperatures, while lower quality composites show

no change.

The available data indicate that the other graphite-epoxy composites have

somewhat lower interlaminar shear strengths, ranging from about 10,000

psi down to about 3000 psi with no systematic temperature dependence being

evident. At 77 K, there appears to be a somewhat higher strength group composed

of Modmor I and II, HMG-25, HMG-50 and HM-S developing strengths on the order

of 7500-10,500 psi, while Samco 360, Thomel 50, Fibralloy 300, GY-70

and Thornel 75 form a somewhat lower group at about 3000-5500 psi. Aside

from the very high strength of the low-modulus HT-S fiber composites,

there is no evidence of a dependence of interlaminar shear strength on fiber

modulus

.

163
Available data on the temperature dependence of inter laminar shear

strengths of boron/ epoxy, boron/aluminijm and PRD 49/epoxy appear on Figure

12. Data reported for experimental Borsic-titanium/ aluminum and Borsic-steel/

13 aluminum hybrids and for experimental CFRP composites of HT-S fiber

58 in polyimide and in epoxy-phenolic matrices are included for comparative

purposes.

The interlaminar shear strength was reported to increase upon cooling for

all of these composites except for the PRD 49/epoxy, which appear^ to be

insensitive to temperature down to 20 K. The highest strength was reported

for the Borsic-titanium/ aluminum hybrid; however, values reported for the

conventional boron/aluminum and SP-272 boron/epoxy composites were almost

13 58 as high, both ranging from about 15,500 psi to 19,000 psi at 77 K. '

The boron/aluminum curve reflects data obtained with both 4.2 mil Borsic and

5.6 mil boron fibers, the highest values at 77 K being reported for the latter.

The HT-S /epoxy-phenolic, the Borsic-steel/ aluminum hybrid and the HT-S/

polyimide composites appear to develop somewhat lower interlaminar shear

strengths, ranging from about 10-14,000 psi at 77 K. The HT-S/epoxy-phenolic

material displayed a particularly large increase in shear strength with decreasing

temperature. A comparatively low value of interlaminar shear strength was

reported for PRD 49/epoxy. The curve for this material reflects results

obtained with both ERLB 4617 and NASA Resin 2 matrices, the latter producing

shear values 40-70% higher than the former?^ '^^

In summary, these data indicate that the best of the HT-S/epoxy CFRP

composites and the conventional boron/ epoxies and boron/ 6061 aluminum

composites are capable of providing equivalent high levels of interlaminar

shear strengths at cryogenic temperatures, the values varying from about

164
14,000 psi at 295 K to about 19,000 psi at 77 K. These values are slightly

higher than those reported in Part I for uniaxial glass-reinforced composites

tested by the short beam method.

Composite Failure Strain

Available data on the temperature dependence of ultimate tensile strain

for advanced composites are minimal; all of that which has been published

13 reflects the work of Hertz, et al. Nevertheless, these data suffice to

illustrate that: (a) the strain at tensile fracture is very low for the advanced

-3

composites, being on the order of 1-9 x 10 in the longitudinal uniaxial

direction and, (b) the strain at tensile fracture is relatively independent of

temperature within the cryogenic range. These data are summarized on Figure

13. These strains are about an order of magnitude lower than those for

glass-reinforced composites, reflecting the negligible fracture strain of the

advanced fibers.

Figure 13 includes data on one CFRP composite (HT-S/X-904) , on two

conventional boron/ aluminum composites (4.2 mil Borsic and 5.6 mil boron,

both in 6061) and on the two Borsic/aluminum hybrids discussed previously

with reference to the other properties. Both longitudinal and transverse ultimate

fracture strains are graphed on Figure 13 for each material.

The data on Figure 13 shows the transverse fracture strain of the HT-S/X-904

-3

CFRP composite to be very low at about 2 x 10 , only about 25% of that

reported for the longitudinal strain. By comparison, the transverse fracture

-3

strain for the commercial HT-S/3002 composite is reported at 7.8 x 10 , only

-3

slightly less than its longitudinal fracture strain of 8.6 x 10 , both room

125

temperature data. Hertz, et al. do not coninent on this low transverse

strain in HT-S/X-904 and, as their data reflect the average of six specimens

taken from panels carefully prepared for design allowable studies, it must be

assumed that these data are valid for the HT-S/X-904 composite.

165
Figure 13 shows that a 5.6 mil boron/ 6061 aluminum composite is expected to

fail at about the same strain in both the longitudinal and transverse directions,

while the 4.2 mil Borsic/6061 composite is expected to fail at a much lower strain

13 when tested in the transverse direction. Hertz, et al. found that the 4.2 mil

Borsic fiber was failing by longitudinal fiber splitting during transverse

testing, while such splitting did not occur in the larger diameter fiber. The

authors concluded that the problem was related to the use of plasma-sprayed tape,

as higher transverse strains were observed in the same type of composite

when diffusion-bonded tape was used to make the test specimens. They concluded

that diffusion-bonded bo ron/ aluminum tape was superior for oriented-ply

126 applications. Kreider has proposed an alternative explanation for longitudinal splitting of the small diameter boron fibers, citing results of diametralcompression tests of individual fibers, which indicate that such splitting reflects

strength anisotropy in the small fibers due to residual stresses retained from

the original fiber manufacturing process and due to residual stresses retained

from the original fiber manufacturing process and from pre-existing flaws in the

fibers. Such anisotropy was not observed in the larger fibers. Fiber splitting

was not directly related to the presence or absence of a silicon carbide coating

on the fibers.*

The Borsic-titanium/aluminum hybrid developed higher strain values in

both directions than did the conventional 4.2 mil Borsic /aluminum, but lower

values than that of the composite reinforced with 5.6 mil boron. The Borsics tee 1/ aluminum hybrid shows clearly the increase in transverse strain

capability due to the presence of stainless wires oriented in this direction.

The longitudinal strain of this hybrid did not appear to be significantly affected

by the transverse reinforcement.

* Manufacturers of boron-aluminum tape report that current fabrication

technology has overcome the problem of fiber splitting.

166
The only data available on the temperature dependence of compressive

strain within the cryogenic temperature range was for the HT-S/X-904

13 CFRP composite, again reflecting the work of Hertz, et al. These data

indicate an increasing compressive strain capability for this composite type

as the temperature is lowered, with a substantially higher strain capability in

the transverse direction. Again, these strain values are substantially lower

125

than those reported for a conventional HT-S/3002 CFRP composite.

Composite Bearing Strength

The only data found in the literature on the temperature dependence

of bearing properties of advanced composites in the cryogenic range were published

13 by Hertz, et al. for boron- and Borsic-reinforced 6061 aluminum, including

the steel and titanium reinforced hybrids. These data, appearing on Figure

15, are for bearing strength in the uniaxial longitudinal direction, i.e.,

the stress required for bearing pull-out, rather than the more conventional

bearing yield strength. The data of Figure 15 refer to material hole- to-free

surface dimensions two and four times the diameter of a No. 10 steel pin

(f^ 0.19 in.). An increase in bearing strength with lowering temperature is

observed for both conventional boron and Borsic-reinforced materials, the

effect being more pronounced in the 4D tests. The Borsic-steel/aluminum

hybrid performed much like the conventional materials, while the Borsictitanium/ aluminum hybrid performed somewhat erratically, although the highest

strengths were developed with this composite.

DYNAMIC MECHANICAL PROPERTIES

Composite Fatigue

The only available data on the temperature dependence of the fatigue

properties of advanced composites were generated during cyclic pressure testing

167
of f ilament-vTOimd pressure vessels. As the composites in these vessels

were crossplied and were subjected to biaxial stressing, the data are of value

only in a comparative sense.

45 Hanson has studied the cyclic fatigue performance of pressure

vessels overwrapped with Thomel 50 graphite fiber using an ERL 2256/ZZL 0820

commercial filament-winding resin. The fatigue life at 77 K was found to be

similar to that at 295 K with 90% of the relevant single-cycle strength being

retained after 10,000 cycles. This compares to a retention of only about 45%

of the single-cycle strength for glass-fiber reinforced vessels. Hanson concluded

that on the basis of specific strength, CFRP vessels would have fatigue performance

superior to that of GFRP vessels after only 80 cycles.

20 Alfring, et al. have reported cryogenic cyclic fatigue data on pressure

vessels filament -wound with 4 mil boron in a Polaris resin matrix.* The results

were again reported to be independent of temperature (down to 20 K) ; however,

the fatigue life of the boron-reinforced vessels was found to be relatively low,

with residual strengths falling to 40-50% of single-cycle values after 10,000

cycles. This is about the same as for glass reinforcement.

These data suggest that graphite may be superior to boron as a reinforcement

fiber for composite structures subject to fatigue at cryogenic temperatures;

however, these data must be considered to be very tentative in view of the

aforementioned tendency for the 4 mil boron fiber to split longitudinally

when subjected to transverse stresses such as are present in biaxiallyloaded pressure vessels. If splitting is adversely affecting the fatigue strength,

the performance of BFRP composites reinforced with 5.6 mil fiber should

be substantially better.

* Polaris resin consists of Epon 828/Epon 1031/NMA/BDMA in proportions

50/50/90/0.55 pbw. Commercial designations are E-787 and 58-68R.

168
Composite Impact Strength

Few data were available on the impact strength of advanced composites

at cryogenic temperatures. The Advanced Composites Design Guide^^^

includes some unpublished data on the Charpy V-notch impact values of

commercial 5505-4 boron/epoxy material in the uniaxial longitudinal direction.

i These data indicate a slight increase in impact strength from 22.8 ft lb in

at 295 K to 27.0 ft lb in" ^ at 20 K. Concomitantly, a slight decrease from 46.3

to 39.5 ft lb in was reported for unnotched specimens over this temperature

range. The notched/ unnotched ratios of about 0.49 at 295 K and about 0.68

at 20 K, indicate some notch sensitivity. Sumner et all have reported

a 12 ft lb in ^ Charpy V-notch strength at 295 K for 25 v/o stainless steel

wire reinforced 2024 aluminum, this value increasing to 18 ft. lb. in

^

at 77 K. Cryogenic impact strength data were not available for the other types

of advanced composites; however, room temperature data indicate that the

CFRP type composites have substantially lower impact strengths than BFRP

type and there is no a priori reason to expect this relationship to change at

cryogenic temperatures. Such impact strengths are substantially lower than

those discussed in Part I for glass-reinforcement where notched values for

-1 4 uniaxial longitudinal composites were seen to range from 67-162 ft lb in at 77 K.

The critical factor has been identified as the stress-strain behavior of the fiber

reinforcement, higher impact values being associated with higher failure

strains of the fiber whenever tests are made in the longitudinal (cross-fiber

^ ^

. , 127,128 fracture) mode.

Impact strengths of uniaxial composites tested in the transverse direction

are essentially matrix controlled. Since composite structures are usually of

crossply construction and subjected to complex stresses, an improvement in

the impact strength of the matrix offers the possibility of an overall improvement

169
of composite toughness. Larsen has reported efforts to improve the impact

properties of crossplied HT-S/epoxy materials at cryogenic temperatures by

addition of elastomeric components to the matrix. Results of this work were

reported in terms of energy density, i.e., the area under the force-deflection

curve using a cleavage-type specimen. This work showed that some improvement

was indeed possible with CBTN modified ERLB 4617 epoxy; however, the benefit

gained was negated by a significant lowering of the room temperature impact

properties and by unpredictable performance at cryogenic temperatures.

Larsen concluded that better overall cryogenic properties would be obtained

by use of the NASA Resin 2 epoxy formulation with CFRP composites.

It appears at this time that substantial improvement of impact strength

of the advanced composites will require development of hybrids, possibly combining glass with the advanced fibers.

THERMAL PROPERTIES

Composite Thermal Expansion and Contraction

Unlike the glass-reinforced composites, some advanced-fiber composites

expand in the fiber direction when cooled to cryogenic temperatures. Available

data for the temperature dependence of dimensional changes in the uniaxial

longitudinal direction are summarized in Figure 16, while comparable data

for the uniaxial transverse direction appear in Figure 17.

The CFRP con5)Osites display a very small longitudinal expansion on

cooling, slightly larger expansions being reported for the high modulus

Thomel 75 and GY-70 fiber composites than for the medium-modulus

Fibralloy 300 and the low-modulus HT-S fiber composites. The slight reversal

in expansion indicated by a dashed line below 77 K for the HT-S composite

is apparently a real effect documented by Larsen and Simon"'"'^ in both

HT-S/NASA Resin 2 and HT-S/4617 Polyblend composites. The PRD 49

170
(Kevlar 49) composites undergo a comparatively large longitudinal expansion

-4

on cooling, reportedly reaching a maximum of about 8-10 x 10 at about 77 K.

Slightly larger expansion was reported with a NASA Resin 2 matrix than with

ERLA 4617. The other advanced composites contracted on cooling. Least

longitudinal contraction was reported for the commercial SP-272 boron/epoxy

product, while the boron/aluminums undergo the largest contractions. The

matrix contraction properties appear to be dominating, in the latter as

relatively little difference is seen among the four variants, including the hybrids

The 5.6 mil composite displays a slightly lower contraction than does the 4 mil

product. The longitudinal thermal contractions of these boron/ aluminum

composites are about 25% of that for unreinforced 6061 aluminum alloy.

Dimensional changes in the transverse direction are strongly influenced

by the matrix, all advanced composites showing contraction in this direction

on cooling. These contractions are large compared to changes observed in the

fiber direction, as witnessed by the necessity to change the ordinate scale

for Figure 17. The data indicate that boron/epoxy composites have the least

transverse contraction of the group, yet even this relatively low value is

three times larger than its longitudinal change. The CFRP composites appear

to undergo about twice the transverse thermal contraction of the BFRP composites, reflecting the lesser constraint provided in this direction by the

small graphite fibers. A slightly higher transverse contraction is reported for

the high-modulus GY-70/X-904 composite than for the low-modulus HT-S/X-904

material; however, such a small difference could easily be accounted for by

variations in fiber volume fraction. The conventional boron/6061 aluminum

composites have the next largest transverse contraction. By comparison

with the contraction of unreinforced 6061 alloy (dashed curve) , it is seen

that the transverse contraction of boron/aluminum is almost completely matrix

171
dominated. No significant differences were reported for 4.2 mil Borsic and 5.6

mil boron fiber composites.

The PRD 49/epoxy composites have by far the largest transverse thermal

contraction. Comparing the temperature dependence of PRD49-I/NASA Resin

2 (curve 8) with that reported by Soffer and Molho^ for MSA Resin 2 alone

(dashed curve), indicates that the matrix is also dominating the transverse

dimensional changes in this composite. The PRD 49 fibers do not offer appreciable

restraint to transverse contraction of the matrix because, as seen in Figure 16,

this fiber expands significantly in the longitudinal direction during cooling

which, of necessity, is accompanied by significant transverse fiber contraction.

Thus, cooling to cryogenic temperatures would be expected to produce sizeable

interfacial shear forces between the fiber and matrix, with lesser forces normal

to the fibers.

Coii^)osite Thermal Conductivity

Few data were available on the temperature dependence of uniaxial longitudinal and uniaxial transverse thermal conductivity of advanced composites.

Those which were available are plotted in Figure 18. In considering these data,

the reader should be aware of the difficulties and potential sources of error

in determining such conductivities in the cryogenic range. The comments

made in Part I of this paper on this subject apply equally well here. Furthermore,

the volume fraction of fiber will influence the conductivities, particularly in

graphite-reinforced composites. For these reasons, the data on Figure 18

should be considered as only indicating trends.

The CFRP composites have relatively high thermal conductivities in the fiber

direction, reflecting relatively high conductivity of the graphite fibers conqjared

to epoxy. Longitudinal conductivity in these materials is highly temperature de172
pendent. The BFRP composites have much lower conductivities than do the CFRP

composites in the fiber direction and show a comparatively small temperature dependence of thermal conductivity. Data were not available for boronreinforced aluminum; however, the thermal conductivity of such materials will

certainly be much higher than that of a polymeric matrix composite. The conductivity should be approximately half that of unreinforced aluminum in a typical

50 volume percent boron/ aluminum composite. Transverse thermal conductivties

are matrix dominated and very low for both the CFRP and BFRP materials.

These data suffice to show that the thermal conductivities of the epoxymatrix advanced composites, particularly boron-reinforced, are low enough to

make them attractive for cryogenic structural components.

Composite Specific Heat

Figure 19 shows the specific heat of the advanced composites to be similar

to those of the glass-reinforced composites discussed in Part I of this paper. Again,

an almost linear temperature dependence of C^ is observed from 295 K to 77 K.

As a group, the values are slightly lower than aluminum for much of the region

between 295 K and 77 K, but substantially higher than that for copper and most

other metals down to 77 K. The temperature dependence of C^ for aluminum,

copper and titanium are included on Figure 19 for comparative purposes.

The highest specific heat at cryogenic temperatures was reported for boron/

aluminum, the aluminum matrix undoubtedly contributing significantly to this

relatively high value. The depicted curve is the average of data reported by

13

Hertz et al. for 4.2 mil Borsic/6061 and 5.6 mil boron/6061. Slightly higher

values were reported at 295 K and 200 K for the larger filament size, falling

to slightly lower at 77 K. The data indicate that the Borsic/ aluminum hybrids

had slightly lower specific heats, probably reflecting the contributions from

173
the steel and titanium reinforcements. The specific heat of the boron/epoxy

composite lies between that of the boron/aluminum and the graphite/epoxy

The specific heat of the latter has undoubtedly been lowered by the graphite

which, as shown by the dashed curve on Figure 19, has itself a relatively

low specific heat.

The data indicate that boron/epoxy has a comparatively high specific heat at

room temperature. This is not unreasonable, as boron has a room temperature

2 -1-1

specific heat of about 12. 9 x 10 J Kg K . No cryogenic specific heat data

were available for boron or epoxy; however, the composite data suggest a rapidly

decreasing value of for boron as temperature decreases, while indicating

that the specific heat of epoxy probably lies between the curves for boron/epoxy

and graphite/epoxy.

COMMENTS ON ADVANCED COMPOSITES

The purpose of this review is to give the reader an understanding of the

present state of knowledge as to the magnitude and temperature dependence of

the properties of advanced fiber composites at cryogenic temperatures. It is

apparent that present knowledge is incomplete in many respects; nevertheless,

the available data justify considerable optimism as to the future applicability

for these materials for structural components operating at cryogenic temperatures.

The cryogenic performance of the commercial state-of-the-art boronreinforced epoxies and aluminum are particularly impressive as they consistently

display excellent tensile and compressive strengths and elastic moduli which

remain constant or iii5)rove as the temperature is lowered. Equally impressive is

the relatively small scatter in the data which is indicative of a high degree of

reliability for components made from boron-reinforced composites. The data suggest

that the commercial products may be used effectively at cryogenic temperatures

without optimization of the composite matrix.

174
Available data on the cryogenic performance of the graphite-reinforced

epoxy class of composites are less convincing. Ultimate tensile strengths are

lower than for boron reinforcement and show a disturbing tendency to decrease

with cooling. Test results are sometimes very unpredictable and contradictory.

Larsen and Simon^^ report diametrically opposed temperature dependence

of the tensile strength for both HT-S/NASA Resin 2 and HT-S/ERLB 4617

polyblend composites when each was tested first as flat uniaxial specimens and

subsequently by the NOL-ring method. Results of this type suggest that much

of the apparent erratic and inconsistent behavior in CFRP composites may be

due to the difficulty of obtaining valid measurements of strength at cryogenic

temperatures with this class of composite. Certainly, the overall performance

of graphite-fiber overwrapped pressure vessels also reported by Larsen and

Simon was much better than could have been predicted from their basic test

data. In contrast to the boron-reinforced materials, data suggest that the CFRP

composites may benefit from matrix optimization. Larsen and Simon recommend

NASA Resin 2 for both uniaxial and crossply CFRP composites on the basis

of their research. This would appear to be a reasonable choice, as this resin

system was developed by Soffer and Molho^ to provide improved performance at

low temperatures. The potential user is cautioned, however, that this resin

has relatively poor properties at room temperature and above. Particular

caution must be taken to properly support components made with this resin

whenever elevated ten5)erature vacuum degassing is required prior to cryogenic

service.

Kevlar 49 (PRD 49) is currently attracting a good deal of attention for

cryogenic applications. This fiber is often considered as a replacement for

glass as it is relatively inexpensive and imparts to a composite a significantly

higher modulus. Based on current knowledge, the NASA Resin 2 epoxy formulation

175
is perhaps the most acceptable matrix material for this fiber; however, the

aforementioned warning on the low room temperature properties must be heeded.

In part I of this paper, the author mentioned the possibility that the

cryogen in which mechanical property measurements are made might be affecting

the results obtained with glass- re info reed polymers. This cautionary note

was based on the recent discovery that the onset of crazing in certain

polymers could be caused by the activity of the liquefied gas (particularly

129 nitrogen) relative to the polymer surface. The possibility exists that

a similar phenomenon is affecting the mechanical properties of the graphiteand PRD 49-reinforced composites. In the few cases where mechanical property

measurements on such composites were carried to below 77 K, the data tended

to become erratic in the same manner as was observed for glass. The ultimate

tensile strength data for the HT-S fiber in Figures 1(a) and 1(b) may be cited

as examples, as may the modulus of PRD 49 in Figure 4(a). It is difficult to

accept such precipitous changes in temperature dependence as inherent

characteristics of the materials.

Much work remains before the behavior of advanced composites under

cryogenic conditions can be said to be thoroughly understood. The follow-on

work suggested in Part I for the glass-reinforced composites applies equally

well to the advanced composites. Following the suggested method of approach,

a program of research is being started at NBS to characterize selected GFRP,

CFRP, BFRP, BFRM (aluminum) and Kevlar 49 composites from 295 K to 4 K both

mechanically and thermally. Hopefully, data derived during the course of this

program will provide firm bases for the understanding of composite performance

at cryogenic temperatures.

176
ACKNOWLEDGMENTS

The author thanks Dr. R. P. Reed for consultation and review of the

completed manuscript. Many thanks are also due the NBS-NOAA library staff

of the Boulder Laboratories for their assistance In retrieving the documents

on which this work depended. This research was supported by the Advanced

Research Projects Agency.

177
APPENDIX

The following materials are referred to in this report:

Fibers

Graphite

HT-S, UthS

HMG-25, HMG-50

Modmor I, II

Thomel 25, 50

Samco 320

Fibralloy 300

GY-70

Boron

Courtaulds Ltd.

Hitco Corp.

Morganite Ltd.

Union Carbide Corp.

Samco Corp

Monsanto Corp.

Celanese Corp.

Borsic

Others

Hamilton Standard Corp.

PRD 49 (Kevlar 49)

Resins

E. I. DuPont de Nemours, Inc.

Epoxies

X-904, X-915

ERL 2256, ERLB 4617

Epon 828, 1031, 58-68R

E-350, E-787

Ciba 8183/137

3002

Others

Skybond 703 (polyimide)

HT-424 (epoxy-phenolic)

Flexibilizers, Coupling Agents, Hardners

ZZL 0820

Empol 1040

Boron Composite Products

SP-272 !. 5505

Miscellaneous Materials

Fiberite Corp.

Union Carbide Plastics Co.

Shell Chemical Corp., Plastics & Resin Div.

U.S. Polymeric Corp.

Hercules Corp.

Monsanto Corp.

American Cyanimid.

Union Carbide Plastics Co.

Emery Industries, Inc.

Minnesota Mining & Manufacturing Co.

Avco Corporation

DSA - dodecenyl succinic anhydride

BDMA - benzyldimethylamine

NMA - nadic methyl anhydride

178
BIBLIOGRAPHY

CONTRACT BIBLIOGRAPHY

I - MECHANICAL PROPERTIES

Contract AF-33(616)-8289

Contractor : Directorate of Materials and Processes, Aeronautical

Systems Division, Air Force Systems Command, WrightPatterson Air Force Base, Ohio.

Research Facility : Narmco Research and Development, San Diego,

California.

1. Brink, N. 0. , "Determination of the Performance of Plastic Laminates

Under Cryogenic Temperatures", ASD-TDR-62-79A , August 1962 (AD

288 944).

1.1 Ibid., "Mechanical Behavior of Reinforced Plastics at Cryogenic

Temperatures", Technical Papers, 20th Annual Technical Conference , Society of Plastics Engineers , Vol. 10, Section XV-2, 1964, pp. 1-19.

1.2 Ibid., "Mechanical Behavior of Reinforced Plastics at Cryogenic

Temperatures", Society of Plastics Engineers Journal , Vol. 20,

1964, pp. 1123-1131.

1.3 Ibid., "Mechanical Behavior of Reinforced Plastics at Cryogenic

Temperatures", Narmco Research and Development Report, Code

No. 105-4, 1964.

2. Chamberlain, D. W. , Lloyd, B. R. and Tennant, R. L. , "Determination

of the Performance of Plastic Laminates at Cryogenic Temperatures",

ASD-TDR-62-794, Part II, March 1964 (N64-24212).

2.1 Chamberlain, D. W. , "Tensile Fatigue Testing at Temperatures

Dovm to 20 K", Advances in Cryogenic Engineering , Vol. 9, 1964,

pp. 131-138.

2.2 Ibid., "Mechanical Properties Testing of Plastic Laminate

Materials Dovm to 20 K", Advances in Cryogenic Engineering , Vol. 10, 1965, pp. 117-125.

Contract NAS 8-11070

Cont ractor : National Aeronautics and Space Administration, George C. Marshall Space Flight Center, Huntsville, Alabama.

Research Facility : Goodyear Aerospace Corporation, Akron, Ohio.

3. Toth, L. W. , Boiler, T. J., Butcher, I. R., Kariotis, A. H. and Yoder,

F. D. , "Program for the Evaluation of Structural Reinforced Plastic

Materials at Cryogenic Temperatures", NASA CR-80061 (Final), August

1966 (N6 7-12051).

179
3.1 Toth, L. W. , "Properties Testing of Reinforced Plastic Laminates

Through the 20 degree K Range" , Technical Papers, 20th Annual

Technical Conference, Society of the Plastics Industry , Section

7-C, 1965, pp. 1-10.

3.2 Toth, L. W. and Kariotis, A. H. , "An Assessment of Test Speci- mens and Test Techniques Useful to the Evaluation of Structural

Reinforced Plastic Materials at Cryogenic Temperatures", Advances

in Cryogenic Engineering , Vol. 10, 1965, pp. 126-133.

3.3 Toth, L. W. , "Properties of Glass-Reinforced Epoxy Through the

20 K Range", Modern Plastics , Vol. 42, August 1965, pp. 123-130.

3.4 Toth, L. W. and Burkley, R. A., "Mechanical Response at Cryogenic Temperatures of Selected Reinforced Plastic Composite

Systems", Goodyear Aerospace Report GER- 13169, Paper No. 16,

Seventieth Annual Meeting of the American Society for Testing and

Materials , June 1967.

3.5 Toth, L. W., Boiler, T. J., Kariotis, A. H. and Yoder, F. D.,

"Program for the Evaluation of Structural Reinforced Plastic

Materials at Cryogenic Temperatures", NASA CR-64005, June

1963 through June 1964 (N65-29724)

Contract NAS 3-6297

Contractor : National Aeronautics and Space Administration, Lewis

Research Center, Cleveland, Ohio.

Research Facility : Aerojet General Corporation, Azuza, California.

4. Lewis, A. and Bush, G. E. , "Improved Cryogenic Resin-Glass Filament

Wound Composites", NASA CR-72163 (Final), March 1967 (N67-31856).

4.1 Lewis, A., Bush, G. E. and Creedon, J., "Improved Cryogenic

Resin/Glass Filament -Wound Composites", NASA Interim Report

CR-54867, May 1966 (N66- 28040).

Contract NAS 6-6287

Contractor : National Aeronautics and Space Administration, Lewis

Research Center, Cleveland, Ohio.

Research Facility : Aerojet General Corporation, Azuza, California.

5. Soffer, L. M. and Molho, R. , "Cryogenic Resins for Glass FilamentWound Composites", NASA CR- 72114 (Final), January 1967 (N67-25076)

5.1 Ibid., "Mechanical Properties of Epoxy Resins and Glass/Epoxy

Conqjosites at Cryogenic Temperatures", Cryogenic Properties of

Polymers , J. L. Koenig, Ed., Marcel Dekker, New York, 1968,

pp. 87-117. (Identical to NASA CR-84451, 1967, (N67-27217).

180
Contract F04701-69-C-0059

Contractor : Space and Missiles Systems Organization, Air Force

Systems Command, Los Angeles Air Force Station,

Los Angeles, California.

Research Facility : The Aerospace Corporation, El Segundo, California.

6. Pepper, R. T. , Rossi, R. E. , Upp, U. W. and Riley, W. E. , "Development of an Aluminum-Graphite Composite", SAMSO-TR-70-301, August

1970 (AD 718 409).

6.1 Pepper, R. T. , Upp, J. W. , Rossi, R. C. and Kendall, E. G.

,

"Aluminum-Graphite Composites", SAMSO-TR-70-174, April

1970 (AD 706 883). (Identical to Metallurgical Transactions , Vol. 2, January 1971, pp. 117-120.)

6.2 Rossi, R. C. , Pepper, R. T. , Upp, J. W. and Riley, W. C.

,

"Development of Aluminum-Graphite Composites", Ceramic

Bulletin , Vol. 50, May 1971, pp. 484-487.

Contract NAS 8-11508

Contractor : National Aeronautics and Space Administration, George

C. Marshall Space Flight Center, Huntsville, Alabama.

Research Facility : Harvey Engineering Laboratories, Torrance,

California.

7. Sumner, E. V. and Davis, L. W. , "Development of Ultrahigh Strength,

Low Density Aluminum Sheet and Plate Composites", NASA CR-85863

(Final), July 1966 (N67-31181).

7.1 Davis, L. W. , "Composites at Low Temperature", Paper No. 15 , Seventieth Annual Meeting of the American Society for Testing

Materials , Jime 1967.

Contract NASA DPR C 10360-B

Contractor : National Aeronautics and Space Administration, Lewis Research Center, Cleveland, Ohio.

Research Facility : Naval Ordnance Laboratory, Silver Springs,

Maryland.

8. Simon, R, A. and Alfring, R. , "Properties of Graphite Fiber Composites

at Cryogenic Temperatures", NASA CR-72652 (NOLTR 69-183), May 1970,

Tasks I and II (AD 746 885).

9. Larsen, J. V. , "Properties of Graphite Fiber Composites at Cryogenic

Temperatures — Effect of Elastomeric Additions to Resin Systems",

NASA CR-72804 (NOLTR 70-195), March 1971, Task III (AD 882 972).

181
10. Larsen, J. V. and Simon, R. A., "Carbon Fiber Composites for Cryogenic

Filament-Wound Vessels", NASA CR- 120899, (NOLTR 71-201), May 1972,

Tasks IV, V and VI, (N7 3- 1155 3)

10.1 Simon, R. A., "Graphite Fiber Composites at Cryogenic Temperatures", Technical Papers, 26th Annual Technical Conference , Society of the Plastics Industry , Section 10-D, 1971, pp. 1-4.

10.2 Larsen, J. W. , "Fracture Energy of CBTN/Epoxy-Carbon Fiber

Composites", Technical Papers, 26th Annual Technical Conference , Society of the Plastics Industry , Section 10-D, 1971, pp. 1-4.

Contract NAS 8-26198

Contractor : George C. Marshall Space Flight Center, Huntsville,

Alabama.

Research Facility : General Dynamics/Convair, San Diego, California.

11. Scheck, W. G. , "Development of Design Data for Graphite Reinforced

Epoxy and polyimide Composites", NASA TN-D2970, Report No. GDC-DBG- 70-005,

Final National Aeronautics and Space Administration, Marshall Space Flight

Center, Alabama, May 1974.

12. Scheck, W. G. , "Development of Design Data for Graphite Reinforced Epoxy

and Polyimide Composites", Report No. GDC-DBG70-005 , General Dynamics

Quarterly Report No. 1, 1970. (See also Maximovich, M. and Scheck, W. G.

Quarterly Report No. 2, 1970).

12.1 Stuckey, J. M. and Scheck, W. G. , "Development of Graphite/

Polyimide Composites", National-Technical Conference, Society

of Aerospace Material and Process Engineers , Vol. 3, 1971, pp.

717-723.

Contract F33615-70-1442

Contractor : Air Force Materials Laboratory, Wright-Patterson Air

Force Base, Ohio.

Research Facility : General Dynamics/Convair, San Diego, California.

13. Hertz, J., Christian, J. L. and Varlas, M. , "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)

13.1 Forest, J. D. , Fujimoto, A. F. and Foelsch, G. F. , "Advanced

Composite Applications for Spacecraft and Missiles, Phase I Final

Report, Volume I: Structural Development", AFML-TR- 71-186, Vol. 1,

March 1972.

13.2 Forest, J. D. and Varlas, M. , "Advanced Composite Applications

for Spacecraft and Missiles, Final Report", AFML-TR- 72 -2 78,

March 1973.

182
13.3 Christian, J. L. and Campbell, M. D. , "Mechanical and Physical

Properties of Several Advanced Metal-Matrix Composite Materials'

Advances in Cryogenic Engineering, Vol. 18, 1973, pp. 175-183.

II THERMOPHYSICAL PROPERTIES

Contract AF-33(657)-9160

Contractor : Air Force Materials Laboratory, Wright-Patterson Air

Force Base, Ohio.

Research Facility : General Dynamics/Astronautics , San Diego,

California

The following reports are in a series entitled "Thermophysical Properties of Plastic Materials and Composites to Liquid Hydrogen Temperature (-423°F)".

14. Haskins, J. F. , Campbell, M. C. , Hertz, J. and Percy, J. L. , MLTDR-64-33, Part I, June 1964 (AD 601 337).

15. Campbell, M. D. , Hertz, J., O'Barr, H. L. and Haskins, J. F., MLTDR-64-33, Part II, March 1965 (X65-18921)

.

16. Campbell, M. D., O'Barr, G. L., Haskins, J. F. and Hertz, J., MLTDR-64-33, Part III, August 1965 (AD 468 155).

16.1 Hertz, J. and Haskins, J. F. , "Thermal Conductivity of Reinforced Plastics at Cryogenic Temperatures" , Advances in Cryogenic Engineering, Vol. 10, 1965, pp. 163-169.

16.2 Campbell, M. D., "Thermal Expansion Characteristics of Some

Plastic Materials and Composites from Room Temperature to

-253°C", Advances in Cryogenic Engineering , Vol. 10, 1965,

pp. 154-162.

16.3 Campbell, M. D., Haskins, J. F., O'Barr, G. L. and Hertz, J. "Thermophysical Properties of Reinforced Plastics at Cryogenic

Temperatures", Journal of Spacecraft , Vol. 3, April 1966, pp.

596-599.

(See also Ref. 13.)

Contract F33615-73-C1388 (Work currently in progress)

Contractor : Air Force Materials Laboratory, Wright-Patterson Air

Force Base, Ohio.

Research Facility : General Dynamics/ Convair, San Diego, California.

183
17. Forest, J. D. and Schaeffer, W. H. , "Advanced Composite Missile and

Space Design Data", General Dynamics Report GDCA-CHB72-001-1,

Progress Report No. 1, October 1972.

18. Forest, J. D., "Advanced Composite Missile and Space Design Data",

General Dynamics Report GDCA-CHB72-001-2, Progress Report No. 2,

January 1973.

Ill - GENERAL BIBLIOGRAPHY

19. Aleck, B., "Fiber glass-Overwrapped 2219-T87 Aluminum Alloy LowPressure Cryogenic Tankage", Society of Aerospace Material and

Process Engineers National Technical Conference, Space Shuttle

Materials , Vol. 3, 1971, pp. 131-134.

20. Alfring, R. J., Morris, E. E. and Landes, R. E., "Cycle-Testing

of Boron Filament -Wound Tanks", NASA CR-72899, National Aeronautics and Space Administration, Lewis Research Center, August

1971 (N71-38023).

21. Barber, J. R. , "Design and Fabrication of Shadow Shield Systems for

Thermal Protection of Cryogenic Propellants" , NASA CR-72595, National

Aeronautics and Space Administration, Lewis Research Center, Cleveland,

Ohio, November 1969 (N70-25098)

22. Baucom, R. M. , "Tensile Behavior of Boron Filament-Reinforced Epoxy

Rings and Belts", NASA TN D-5053, Langley Research Center, Hampton,

Virginia, March 1969 (N69-19918).

23. Benton, W. , Carr, R. , Cohen, A., Gustafson, G., Lankton, C. and

Zeldin, B., "Propellant Storability in Space", RPL-TDR-64-75 (Final),

Air Force Systems Command, Edwards Air Force Base, California,

June 1964 (AD 603 215).

24. Brechna, H. , "Superconducting Magnets for High Energy Physics Applications", Proceedings 1st International Cryogenic Conference, Tokyo

and Kyoto, Japan, April 1967 , Heywood Temple Industrial Publishers,

Ltd., London, 1968, pp. 119-129 (CFSCI N67-36009)

25. Bullard, B. R. , "Cryogenic Tank Support Evaluation", NASA CR-72546,

NASA Lewis Research Center, Cleveland, Ohio, December 1969 (N70-13085).

26. Campbell, M. D. , "Development of Thermal Expansion Capabilities and

the Investigation of Expansion Characteristics of Space Vehicle Materials",

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184
27. Campbell, M. D. , "Development of the Thermal Expansion Capabilities

and the Investigation of the Thermal Expansion Characteristics of Space

Vehicle Materials (II)", General Dynamics/Astronautics Report ERRAN-450, 26 December 1963.

28. Caren, R. P., Coston, R. M. , Holmes, A. M. C. and Dubus, F. , "LowTemperature Tensile, Thermal Contraction, and Gaseous Hydrogen

Permeability Data on Hydrogen-Vapor Barrier Materials", Advances in

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29. Chiao, T. T. and Moore, R. L. , "Tensile Properties of PRD-49 Fiber

in Epoxy Matrix", Journal of Composite Materials , Vol. 6, October 1972,

pp. 547-551.

30. Cooper, G. A. and Sillwood, J. M. , "Multiple Fracture in a Steel Reinforced Epoxy Resin Composite", Journal of Materials Science , Vol. 7,

1972, pp. 325-333.

31. Darwish, F. and Tetelman, A. S. , "Mechanical Behavior of Si02-Epoxy

Composite", Conference Proceedings No. 63, Advisory Group for Aerospace Research and Development, Symposium on Composite Materials , Paper No. 9, Paris, April 1970 (Hartford House, London)

.

32. Davis, J. G. and Zender, G. W., "Mechanical Behavior of Carbon Fiber

Reinforced-Epoxy Composites", 12th National Symposium, Society of

Aerospace Material and Process Engineers , Vol. 12, 1967, Section AC-10.

33. Dervy, A. J. , "Reinforced Plastics of High Strength/Weight Ratio for

Space Applications" , Technical Papers, 17th Annual Technical Conference , Society of the Plastics Industry , Section 7-D, 1962, pp. 1-10.

34. Fontana, M. G., Bishop, S. M. and Spretnak, J. W. , "Investigation of

Mechanical Properties and Physical Metallurgy of Aircraft Alloys at Very

Low Temperatures, Part 5 - Mechanical Properties of Metals and a Plastic Laminate at Low Temperatures", AF Technical Report 5662,

Part 5, Materials Laboratory, Wright-Patterson Air Force Base, Ohio,

December 1953 (AD 27726).

35. Freeman, S. M., "Properties of Vapor Barriers, Adhesives and Foams

at Cryogenic and Elevated Temperatures", Lockheed Aircraft Corporation

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36. Freeman, W. T. and Campbell, M. D., "Thermal Expansion Characteristics of Graphite Reinforced Composite Materials", Composite Materials : Testing and Design (Second Conference) , ASTM STP 497, American Society

for Testing and Materials, 1972, pp. 121-142.

185
37. Funk, C. W. and Dixon, C. E., "Cryogenic Radiation Damage in

Structural Polymers", Transactions of the American Nuclear Society , Vol. 9, 1966, pp. 406.

38. Gille, J. P., "Development of Advanced Materials for Integrated Tank

Insulation System for the Long Term Storage of Cryogens in Space",

NASA CR-102570 (Final), National Aeronautics and Space Administration,

Huntsville, Alabama, September 1969 (N70-23348)

39. Gleich, D., "Development of a Filament-Overwrapped Cryoformed Metal

Pressure Vessel", NASA CR-72753, National Aeronautics and Space

Administration, Lewis Research Center, Cleveland, Ohio, January 1971

(N71- 22401) .

40. Gray, P. D., Cornelius, G. K. , O'Donnell, J. D. , and Howard, W. W.

"Rockets in Space Environment, Volume 1: The Experimental Program",

RTD-TDR-63-1050, Aerojet General Corporation, February 1963 (N63-

20999).

41. Greer, F. , "Flexural Properties of Conolon 506 at Room Temperature,

-320°F and -423°F", Convair /Astronautics Report 55E 522, June 1961

(AD 677 565).

42. Hale, D. V., "Study of Thermal Conductivity Requirements: MSFC 20-

Inch and 105-Inch Cryogenic Tank Analyses", NASA CR-61288, National

Aeronautics and Space Administration, Marshall Space Flight Center,

Alabama, June 1969 (N69-35811)

43. Hall, J., "Cryogenic Tensile Tests - Epoxy Fiberglas" , Douglas Aircraft

Company Report MP 1348, September 1961.

44. Hanson, M. P., "Effects of Temperature and Creep Characteristics of

PRD-49 Fiber/Epoxy Conqjosites" , NASA TN D-7120, National Aeronautics and Space Administration, Lewis Research Center, Cleveland,

Ohio, November 1972 (N73- 12607)

45. Hanson, M. P. , "Tensile and Cyclic Fatigue Properties of Graphite

Filament-Wound Pressure Vessels at Ambient and Cryogenic Temperatures", NASA TN D-5354, National Aeronautics and Space Administration,

Lewis Research Center, Cleveland, Ohio, July 1969 (N69- 31300) . (Identical to SAMPE Vol. 15, pp. 249-256).

46. Hanson, M. P., Richards, H. T. and Hickel, R. 0., "Preliminary Investigation 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,

March 1965.

186
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, February 1968. (Identical to NASA

TM X-52350, 1967.)

48. Hanson, M. P., "Static and Dynamic Fatigue Behavior of Glass FilamentWound pressure Vessels at Ambient and Cryogenic Temperatures", NASA

TN D-5807, National Aeronautics and Space Administration, Lewis Research

Center, Cleveland, Ohio, May 1970 (CFSTI-CSCL-20 K)

.

49. Haskins, J. F. and Hertz, J., "Thermal Conductivity Testing of Coast

F-224-6 Phenolic-Fiberglass Laminate", General Dynamics/ Convair

Report No. AR-592-1-482 , July 1963.

50. Haskins, J. F. and Hurlich, A., "Measured Values for the Coefficients

of Linear Expansion of Plycel 420 and Conolon 506 at Low Temperatures",

Convair/Astronautics Report MRG-154 , May 1960.

51. Haylett, J. W., Rottmayer, E. and Butcher, I., "Advanced Composite

Material Study for Millimeter Wavelength Antennas", Technical Report

AFML-TR-71-205, Vol. 1, Air Force Materials Laboratory, WrightPatterson Air Force Base, Ohio, October 1971 (AD 893 368).

52. Haylett, C. E. , "Advanced Composite Material Study for Millimeter

Wavelength Antennas, Volume II, Environmental Tests", AFML-TR-71-

205-Vol. 2, Air Force Materials Laboratory, Wright-Patterson Air

Force Base, Ohio, October 1971 (AD 893 358 L)

.

53. Herring, H. W., Baucom, R. M. and Pride, R. A., "Research on Boron

Filaments and Boron Reinforced Composites", 10th National Symposium , Society of Aerospace Material and Process Engineers , Vol. 10, 1966,

pp. B-21 to B-34.

54. Hertz, J., "Tensile Testing of Conolon 506 at Room and Sub-Zero Temperatures", Convair/Astronautics Report MRG-120, December 1959.

55. Hertz, J., "Tensile Testing of Adlock 851, Adlock PG-LA and Adlock

EG-11A-81A from -423°F to 78°F", Convair/Astronautics Report MRG

237, June 1961.

56. Hertz, J., "Investigation of Potential Low Temperature Insulations",

General Dynamics/ Astronautics Report GS/A-ERR-AN-668 , December 1964.

57. Hertz, J. , "The Effect of Cryogenic Temperatures on the Mechanical

Properties of Reinforced Plastic Laminates", General Dynamics Report

AR-592-1-415, March 1963 (AD 405 170).

187
58. Hertz, J., "Investigation into the High-Temperature Strength Degradation of Fiber-Reinforced Resin Composite During Ambient Aging",

General Dynamic s/Convair Report No. GDCA-DBG7 3-005 (Final Contract

NAS 8-27435), June 1973.

59. Hoggatt, J. T., "Development of Cryogenic PRD-49-1 Filament -Wound

Tanks", NASA CR-120835, National Aeronautics and Space Administration,

Lewis Research Center, Cleveland, Ohio, December 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 Shuttle

Materials , Vol. 3, 1971, pp. 157-167.

61. Hust, J. G. , "Thermal Conductivity of an Epoxy-Fiberglass Laminate",

National Bureau of Standards Cryogenic Division Unpublished Laboratory

Note 73-1, January 1973.

62. Johnston, H. L. and Brooks, H. E. , "Impact Strength of Various Metals

at Temperatures dovm to 20° Absolute", Ohio State University Cryogenic

Laboratory Report TR 264-17, May 1952.

63. Keller, C. W. , "Fiberglass Supports for Cryogenic Tanks", NASA CR120937 (Final) , National Aeronautics and Space Administration, Lewis

Research Center, Cleveland, Ohio, October 1972 •(N72-33564)

64. Kerlin, E. E. and 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 Space Administration, George C. Marshall Space Flight

Center, Huntsville, Alabama, July 1966 (N66-35963)

65. Kerlin, E. E. and Smith, E. T., "Measured Effects of the Various Combinations of Nuclear Radiation, Vacuum and Cryotemperatures on Engineering Materials: Annual Report", NASA CR-58830, National Aeronautics and Space Administration, George C. Marshall Space Flight

Center, Huntsville, Alabama, May 1964 (N64-33043)

66. Keys, R. D. , Kiefer, T. F, and Schwartzberg, F. R., "Cryogenic

Properties of High-Strength Glass-Reinforced Plastics", Advances in

Cryogenic Engineering , Vol. 11, 1966, pp. 470-477.

67. Krause, D. R. , "Development of Lightweight Material Composites to

Insulate Cryogenic Tanks for 30-Day Storage in Outer Space", NASA CR123797, National Aeronautics and Space Administration, George C.

Marshall Space Flight Center, Huntsville, Alabama, June 1972 (N72-

30495).

188
68. Krause, D. R. , Fredrickson, G. 0. and Klevatt, P. L., "Effects 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,

pp. 639-643.

69. Lantz, R. B., "Materials for Filament Wound Cryogenic Pressure

Vessels", 6th National Symposium, Society of Aerospace Materials

and Process Engineers , Vol. 2, Engineering Paper No. 1750.

70. Lavengood, R. E. and Anderson, R. M., "Matrix Properties Controlling

Torsional Fatigue Life of Fiber Reinforced Composites", Technical

Papers, 24th Annual Technical Conference, Society of the Plastics

Industry , Section 11-E, 1969, pp. 1-7.

71. Levin, V. A., Naumenkov, P. G. and Shchitov, M. V., "Some Properties

of Plastics at Low Temperatures", Plasticheskia Massy , Vol. 11, 1966,

pp. 64-65.

72. Luikov, A. V., Vasiliev, L. L. and Shashkov, A. G., "A Method for the

Simultaneous Determination of all Thermal Properties of Poor Heat

Conductors Over the Temperature Range 80 to 500 K", Proceedings , Third American Society of Mechanical Engineers Symposium , Purdue

University, March 1965, pp. 314-319.

73. Lyon, D. N. and Parrish, W. R., "Low Temperature Thermal Conductivities of Two High Compressive Strength Materials", Cryogenics , Vol. 7, No. 1, 1967.

74. Maher, L. E., "Some Problems Arising from the Use of HydrogenFuelled Propulsion Systems", Journal of the Royal Aeronautical Society , Vol. 68, November 1964, pp. 765-771.

75. McKannon, E. C. and Cause, R. L., "Effects of Nuclear Radiation

and Cryogenic Temperatures on Non-Metallic Engineering Materials",

Journal of Spacecraft , Vol. 2, August 1965, pp. 558-564.

76. Morris, E. E. , "Glass-Fiber-Reinforced Metallic Tanks for Cryogenic

S e rvice " , 12th National Symposium, Society of Aerospace Materials and

Process Engineers , Vol. 12, October 1967, Section AS-4. (Also NASA

CR-72224.)

77. Morris, E. E., "The Performance of Glass-Filament -Wound Pressure

Vessels with Metal Liners at Cryogenic Temperatures", Journal of

Materials . Vol. 4, December 1969, pp. 970-1004.

189
78. Morris, E. E. and Alfring, R. J., "Cryogenic Boron-Filament-Wound

Pressure Vessels", Composite Materials; Testing and Design, ASTM

STP 460 . American Society for Testing and Materials, 1969, pp. 430-443.

79. Morris, E. E. and Landes, R. E., "Cryogenic Glass-Filament-Wound

Tank Evaluation", NASA CR~72948 (Final), National Aeronautics and

Space Administration, NASA Lewis Research Center, Cleveland, Ohio,

July 1971 (N72-14696).

80. Mowers, R. E. , Leib, J. H. and Sherman, S., "Program of Testing

Nonmetallic Materials at Cryogenic Temperatures", Rocketdyne Corporation Report R-3498, Rocket Propulsion Laboratories, Edwards,

California, December 1962 (AD 294 772).

81. Nadler, M. A., Yoshino, S. Y. and Darms, F. J., "Boron/Epoxy

Support Strut for Non-Integral Cryogenic Tankage", North American

Rockwell Space Division Report SD 68-99501, February 25, 1969. (See

also 15th National Symposium SAMPE April 1969 and North American

Rockwell Report SD 995 2, 1968.)

82. Nelson, L. F., "Compressive Strength of Conolon 506 at +75°F and

-320°F", Convair/Astronautics Report No. 27E 1336, January 1962.

83. Nelson, L. R. , "Mechanical Properties of Adlock 851 at Room Temperature, 1000°, -320° and -423°F", Convair/Astronautics Report No. 55E

812, July 1961.

84. Nelson, P. T. and Archer, J. S., "Graphite Reinforced Plastic EHF

Antenna", TRW Systems Group, Redondo Beach, California Report No.

99900-7128-RO-ll, December 1969.

85. Patten, P. M. , "Internal Insulation Liner Alteration", Douglas Aircraft

Company Report No. SM 45975, August 1964.

86. Perkins-Elmer Optical Group, Norwalk, Connecticut, Work-in-Progress

on Contract No. F33615-72-C-2033, Air Force Systems Command, WrightPatterson Air Force Base, Ohio.

87. Pink, E. and Campbell, J. D., "The Effect of Strain Rate and Temperature on the Deformation Behavior of Reinforced and Unreinforced Epoxy

Resin", Oxford University Department of Engineering Report No. 1040/72,

Oxford, England, July 1972 (N73-10568)

88. Pirgon, 0., Wostenholm, G. H. and Yates, B. , "Thermal Expansion at

Elevated Temperatures, IV. Carbon-Fibre Composites", Journal of

Physics D: Applied Physics , Vol. 6, 1973, pp. 309-321.

190
89. Pride, R. A., Stein, B. A. and 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 Management Conference , Washington,

DC, 1968, Section 17-c, pp. 1-8.

90. Ratcliffe, E. H., "Thermal Conductivities of Plastics with Glass,

Asbestos and Cellulosic Fiber Reinforcements", Applied Material Research , Vol. 5, 1966, pp. 200-201.

91. Roseland, L. M. , "Materials for Cryogenic Usage", Technical Papers , 21st Annual Technical Conference, Society of the Plastics Industry , 1966,

Section 4-C, pp. 1-6.

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. SM-48409, January 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, June 1959 (AD 830 230).

94. Sanders, R. H. and Weleff, W. , "Final Report on GTR-17 Effects of

Radiation on Organic Materials Irradiated in Liquid Hydrogen", AerojetGeneral Corporation Report No. RN-S-0327, March 1967.

95. Sanger, M. J., Molho, R. and Howard, W. W., "Exploratory Evaluation

of Filament -Wound Composites for Tankage of Rocket Oxidizers and

Fuels", AFML-TR-65-381, Air Force Materials Laboratory, WrightPatterson Air Force Base, Ohio, January 1966 (AD 477 455).

96. Sanger, M. J. and 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 , 1967,

Section AS-7.

97. Sewell, J. J. and Kuno, J. K. , "Aerospace Use of Plastic Hardware and

Thermal Insulation", Technical Papers, 17th Annual Technical Conference , Society of the Plastics Industry , 1962, Section 7-A, pp. 1-14.

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, November 1964

(N65-10775).

191
Soltysiak, D. J. and 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 , 1967, Section 14-E, pp. 1-14.

Speare, J. C. , "Preliminary Sizing of Filament -Wound RNS Tanks",

Report No. TOR-0066 (5 759-07)-13 , Space and Missile Systems Organization, Air Force Systems Command, Los Angeles Air Force

Station, Los Angeles, California, June 1970 (AD 872 626).

Steinhauer, R. A., "Linear Thermal Expansion of 828CL 181 Cloth

Laminate", Douglas Aircraft Company Report No. MP 11,979, August

1961.

Stinnett, W. D., "Cryogenic Tensile Properties of Selected Materials

NASA CR-71751, AEC-NASA Space Nuclear Propulsion Office, Report

No. 2712, January 1964 (N66-22816)

Suezawa, Y., Hojo, H. and Nakamura, K. , "Impact Characteristics

of Fiberglass Reinforced Plastics at Low Temperatures", Kagaku

Kogaku (Chemical Engineering, Japan) , Vol. 33, 1969, pp. 1051-1059.

Toth, J. M., "Barrier Films for Filament -Wound Fiberglass Cryogenic

Vessels", Advances in Cryogenic Engineering , Vol. 1, 1964, pp. 537-

544.

Toth, J. M. and 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, May 1967 (N67-25070).

Toth, J. M. , Sherman, W. C. and Soltysiak, D. J., "Investigation of

Smooth-Bonded Metal Liners for Glass Fiber Filament-Wound Pressure

Vessels", Douglas Missile and Space Systems Division Report No. SM49384, Quarterly Report No. 3, Contract No. NAS 3 6293, NASA Lewis

Research Center, Cleveland, Ohio, April 1966.

Toth, J. M. , Sherman, W. C. and 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).

Toth, J. M. and Barber, J. R., "Structural Properties of Glass-Fiber

Filament-Wound Cryogenic Pressure Vessels", Advances in Ciryogenic

Engineering . Vol. 10, 1965, pp. 134-144.

192
Voloshenko-Klimovitskii, Yu. Ya. , Belyaev, Yu. A., L'vof, B. S.

and 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 , Vol. 6, 1964, pp. 39-40.

Voloshenko-Klimovitskii, Yu. Ya. , Belyaev, Yu. A. and Korenkov,

Yu. A. , "Impact Tensile Tests on Glass-Fibre Reinforced Plastics

at Normal and Low Temperatures", Plasticheski Massy , No. 5, 1963,

pp. 51-54.

Watson, J. F. , Christian, J. L. and Hertz, J., "Selection of Materials

for Cryogenic Applications in Missiles and Aerospace Vehicles",

Convair/Astronautics Report No. MRG 132-1, February, 1960.

Weleff, W. , "Effect of Nuclear Radiation and Liquid Hydrogen on Mechanical Properties of Three Phenolic Materials", Advances in

Cryogenic Engineering , Vol. 11, 1966, pp. 486-491.

Weleff, W. , "Final Report, GTR-16 Radiation Effects Test on Structural

Materials at -423°F", Aerojet-General Corporation Report No. RN-S0290, November 1966.

IV - HANDBOOKS AND REVIEWS

Coston, R. M. , "Handbook of Thermal Design Data for Multilayer

Insulation Systems", LMSC-A847882, Vol. II (Final), George C.

Marshall Space Flight Center, Huntsville, Alabama, June 1967 (N67-

34910) .

Hertz, J., "The Effect of Cryogenic Temperatures on the Mechanical

Properties of Reinforced Plastic Laminates", Society of Plastics

Engineers Journal , Vol. 21, February 1965, pp. 181-189.

Hertz, J. and Knowles, D., "Survey of Thermal Properties of Selected

Materials", General Dynamics /Convair Report AAL-65-008 (AR-504-

1-553), February 1965 (N65-31775).

Jurevic, W. G. and Rittenhouse, J. B., "Structural Applications

Handbook", AFML TR-67-332, Air Force Materials Laboratory,

Wright-Patterson Air Force Base, Ohio, August 1968 (AD 804 585).

Lackman, L. M. , Arvin, G. H. , et al. , "Advanced Composite Design

Guide, Third Edition, Volume IV: Materials", Air Force Materials

Laboratory, Wright -Patters on Air Force Base, Ohio, January 1973.

193
Landrock, A. H., "Properties of Plastics and Related Materials at

Cryogenic Temperatures", Plastec Report No. 20, Plastics Technical

Evaluation Center, Picatinny Arsenal, Dover, New Jersey, July 1965

(AD 469 126).

Maximovich, M. and Scheck, W. G. , "Data Summary and Reference

File for Graphite and Boron Reinforced Composite Materials", General

Dynamics/Convair Report No. GDCA-DBG71-006 , 1971 (Contract NAS

8-26198, George C. Marshall Space Flight Center, Huntsville, Alabama).

Nored, D. L. , Hennings, G. , Sinclair, D. H., Smith, G. T., Smolak,

G. R. and Stofan, A. J., "Storage and Handling of Cryogenic Fluids",

NASA Special Publication SP-5053, Proceedings of Converence on Selected Technology for the Petroleum Industry, Lewis Research Center,

Cleveland, Ohio, December 1965 (N66-33674)

"Plastics for Aerospace Vehicles, Part 1. Reinforced Plastics", MILHDBK-17A, Department of Defense, Washington, DC, January 1971.

Rittenhouse, J. B. and Singletary, J. B., "Space Materials Handbook",

Third Edition, NASA Special Publication SP-3051, National Aeronautics

and Space Administration, Washington, DC, 1969 (Limited publication

as AFML-TR-68-205)

Schwartzberg, F. R., Hertzog, 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, July 1970 (AD 713 620).

V - MISCELLANEOUS REFERENCES

"Advanced Composites Design Guide", Third Edition, Volume IV,

Materials, Air Force Materials Laboratory, Wright-Patterson Air

Force Base, Ohio, January 1973.

Kreider, K. G. and Prewo, K. M., "The Transverse Strength of Boron

Fibers", Composite Materials, Testing and Design , ASTM STP 497, 1972,

pp. 539-550.

Arp, v., Wilson, J. H. , Windich, L. and Sikora, P., "Thermal Expansion

Some Engineering Materials from 20 to 293 K" , Cryogenics , June 1962,

p. 230.

Novak, R. C. and DeCrescente, M. D., "Impact Behavior of Unidirectional

Resin Matrix Composites Tested in the Fiber Direction", Composite

Materials, Testing-and-Design, ASTM STP 497, 1972, pp. 311-323.

Kastelic, J. R. , Hiltner, A. and Baer, E., "Crazing, Yielding and

Fracture in Polycarbonate and Polyethelene Terepthalate at Low Temperatures", Journal of Macro Molecular Science-Physics , B7(4), 1973,

pp. 679-703.

194
I ro O

Oj sQ >-4

Ol t— ( « J- » m !

I O J- • J sD ^ .

\0

rvj " o " -

, J

l - -

- - r— so

iTN sD ^ ^ m iTvsO ^

OJ ro m m iN

~ o

X- o

CVJ J- -

OJ •• - - -

. -co -

rvj irv OJ \0

I •

C\J J -

J- - -

* " Ol ir\ \D

i/v ^ no ^ - *

I I • ^ -H no l/^ k/N ^ ro in

OJ r-H

n-1 - ^ rn .

I I J-

-I cn OJ

* OJ o

no i/N no rn l/^ JnO\8

^ rj ^

w CO no

- nOsO J no

SO J- O

195
OJ

CO

0< o

CO

ij ^

•H

D ^

CM vo

+>

Q O

oc p.

ft, :o

33 ',

OJ CO

o fnoD ^ ^ LTN o oT r-^ rH ^ -dO. C\J OJ O

-IN

o

Iao

o

Ias

4J •

5 r^ -a

3s

U £ Li

o rs H ^ —

>> -H

- --^ o

cd a)

r-l t,

196
Composite System

s-glass/ epoxy

Kevlar 49/epoxy

boron/epoxy

graphite / epoxy

boron/ altiminum (4 mil)

TABLE 1

TYPICAL COMPOSITE DENSITIES

Fiber/Re sin Ratio

60-67

60-65

55

55-60

50

-3 -3 -4

Density, lb in (kg m xlO )

0.068-0.074 (1.87-2.04)

0.047-0.050 (1.3-1.38)

0.070-0.074 (1.93-2.04)

0.050-0.055 (1.38-1.52)

0.10 (2.62)

TABLE 2

COMPARISON OF RANKING OF SPECIFIC UNIAXIAL GRAPHITE /EPOXY COMPOSITES

BY LONGITUDINAL FLEXURE STRENGTH AND BY

LONGITUDINAL TENSILE STRENGTH (77 K)

Composite Type

fu

KSI Rank g"", KSI Rank

HT-S/X-904

HT-S/E-350

HT-S/X-915

HT-S/Ciba 8183

GY-70/X-904

Thomel 50 /Re sin 2

HMG-25/Resin 2 Thomel 50/ERL 2256

178

172

167

146

117

58

55

54

147

87

103

45

92

91

71

93

197
LIST OF FIGURES

Page

tu Figure 1 Ultimate Tensile Strength, a , of Graphite /Epoxy Composites . . 199

Figure 2 Ultimate Tensile Strength, a^", of Miscellaneous Advanced

Composites 200

Figure 3 Initial Tensile Modulus, E^, of Graphite/Epoxy Composites . . . 201

Figure 4 Initial Tensile Modulus, E^, of Miscellaneous Advanced

Composites 202

Figure 5 Ultimate Flexural Strength, o^", of Graphite/Epoxy Composites. . 203

Figure 6 Ultimate Flexural Strength, a^", of Miscellaneous Advanced

Composites 204

Figure 7 Flexural Modulus, E^, of Advanced Composites 204

cu Figure 8 Ultimate Compressive Strength, a , of Advanced Composites . . . 205

Figure 9 Compressive Modulus, E^, of Advanced Composites 206

Figure 10 Histograms Illustrating the Reported Range; of Longitudinal

si

Interlaminar Shear Strength, a , as reported for HT-S/X-904

Graphite/Epoxy at 77 K (a) and the Reported Changes in

Interlaminar Shear Strength in the Same Composite Upon

Cooling From 295 K to 77 K 207

s i Figure 11 Interlaminar Shear Strength, o , of Graphite/Epoxy Composites

Data Reflect Highest Values Reported (See Text for Details)

Flat Short-beam Test Data except as noted 208

si Figure 12 Interlaminar Shear Strength, o , of Miscellaneous Advanced

Composites. Flat Short-beam Test Data Except Where Noted .... 209

Figure 13 Ultimate Tensile Strain, e^^, of Advanced Composites 210

CU Figure 14 Ultimate Compressive Strain, e , of Graphite/Epoxy Composites. . 211

Figure 15 Ultimate Bearing Strength, a^", of Boron/Aluminum Composites,

Uniaxial Longitudinal Direction 212

Figure 16 Longitudinal Thermal Expansion, AL/L, of Uniaxial Advanced

Composites 213

Figure 17 Transverse Thermal Expansion, AL/L, of Uniaxial Advanced

Composites 214

Figure 18 Thermal Conductivities, X, of Advanced Composites 215

Figure 19 Specific Heat, C , of Advanced Composites 216

198
xlO~

170

160

150

140

130

120

no

100

90

80

70

60

50

9

8

7

6

5

4

3

2

1

(1) Thome! 75 {Ref.51) (9) HT-S (Ref .9,10,13,13.2,58]

.(2) GY-70 (Ref. 13) (10) Modmor II (Ref. 8)

(3) Thornel 50 (Ref ,8,10)(11) HMG-25 (Ref. 8)

(4) HMG-50 (Ref. 10)

"(5) HM-S (Ref. 10) ^x^(l)*

(6) Modmor I (Ref. 8)

-(7) Samco 320 (Ref. 8)

^ . . (8) Thornel 25 (Ref. 47)

- UNIAXIAL LONGITUDINAL

(1) Thornel 75/ ERLB 4617 (Ref.51)

(2a) HT-S/NASA Resin 2 (Ref. 10) % (2b) HT-S/4617 Polyblend (Ref. 10) \ (2c) HT-S/X-904)(Ref .13)

\ \ (2a)

(b)

UNIAXIAL TRANSVERSE

XlO'

50 100 150 200

TEMPERATURE, K

250 300

FIGURE 1 - Ultimate Tensile Strength, a^^, of Graphi te-Epoxy Composites.

199
COa.

170

160

150

140

130

120

^ 110

100

90

80

70

60

50

28

24

20

16

12

(1) PRD-49 I/Epoxy (Ref. 59,60)

(2) Boron/Epoxy (Ref . 13,22,47,53,58,81)

(3) HT-S/Polymide (Ref. 13)

(4) HT-S/Epoxy-Phenolic (Ref. 58)

(5) Boron/Aluminum (6061) (Ref . 13)

(6) Borsic-Steel/Aluminum (6061)

(Ref. 13)

(7)

(7) Borsic- Titanium/Aluminum (6061) (Ref . 13)

(8) Stainless Steel /Aluminum (2024)(Ref .7,7. 1)

(b)

UNIAXIAL TRANSVERSE

(3)

.(1) Boron/Epoxy (Ref. 13,81)

(2) Boron/ Alumuinum (6061) (Ref . 13)

"(3) Borsic-Titanium/Aluminum (Ref. 13)

0 50 250

xlO

13

12

11

10

8

I E

I

2.0

1.8

1.6

1.4

1.2

1.0

0.8

0.6

100 150 200 300

TEMPERATURE, K

FIGURE 2 - Ultimate Tensile Strength, a^^, of Miscellaneous Advanced

Composi tes

200
xlO'

50

45

40

35

ao

25

20

15

10

<10*

(1) Thornel 75 (Ref.51)

(2) Thornel 50 (Ref.8)

"(3) Modmor I (Ref.8)

(4) Samco 320 (Ref.8)

(5) HT-S (Ref.9,10,13)

(6) HMG-25 (Ref.8)

xlO

35

- 30

- 25

10

High Modulus

Medium Modulus

Low Modulus

(2)

(a) UNIAXIAL LONGITUDINAL

(6)

Minimal Data

1

(5)

J.

- 20

CM 15

- 10

(1) Thornel 75/ERLB 4617 (Ref.51) (2b) HT-S/4617 Polyblend

-(2a) HT-S/NASA Resin 2 (Ref.lO) (Ref.lO)

^ • (2c) HT-S/X-904 (Ref.l3)

"""""^--^

(2c)

"^--^

(1) ^

(b) UNIAXIAL TRANSVERSE

I I

'*^---..^(2a)

-^^^^£2^b)

xlO12

10

8

6

4

2

0

0 50 100 150 200 250 300

TEMPERATURE, K

FIGURE 3 - Initial Tensile Modulus, E^, of Graphi te-Epoxy Composites

201
xlO

40

35

30

25

20

15

25

20

15

10

5

0

xlO^

(a) UNIAXIAL LONGITUDINAL

(1) Boron/Epoxy (Ref. 81)

(2) PRO 49-I/Epoxy (Ref. 59,60)

(3) Boron/Aluminum (6061) (Ref . 13)

(4) Borsic-Titanium/Aluminum (6061) (Ref . 13)

(5) Borsic-Steel /Aluminum (6061) (Ref . 13)

Minimal Data

(b) UNIAXIAL TRANSVERSE

- (1) Boron/Epoxy (SP-272. Narmco 5505) (Ref .81)

(2) Boron/Aluminum (6061 ) (Ref . 13) - (3) Borsic-Titanium/Aluminum (6061) (Ref . 13)

(4) Borsic-Steel/Aluminum (6061) (Ref . 13)

xlO

26

24

22

20

18

16

14

12

XlO

16

- 12

- 8

- 4

50 100 150 200

TEMPERATURE, K 250 300

FIGURE 4 - Initial Tensile Modulus, E^, of Miscellaneous Advanced

Composites.

2 02
xlO200 r

175 -

150

125

If) a.

o 75 -

50 -

(a) UNIAXIAL LONGITUDINAL

1—•**'^

(1) GY-70 (Ref. 11,13)

(2) Thornel 50 (Ref. 8)

(3) Courtaulds HM (Ref. 8)

(4) Fibralloy 300 (Ref. 13)

(5) HT- S (Ref. 11, 12, 13, 58)

(6) Modmor II (Ref. 13)

(7) HMG-25 (Ref. 8)

High Modulus

ICQ - — Medium Modulus

Low Modulus

(1)

(4)^

Minimal Data

(7)^

(2) (3)^

xlO

14

8

12

10

xlO20

15

Modmor 11/1004

HT-S/1004 ^-^

HT-S/BXP-2401-

HT-S/X-904-^

10 -

HT-S/3002'

GY-70/ 1004

GY-70/R-350A^

(b) UNIAXIAL TRANSVERSE

^iri.^—.

, ^y^^^ (Ref.ll,13)~

(Ref. 13)

6Y-70/X-904 (Ref. 11,13)

I

xlO

12

10

8

6

50 100 150 200

TEMPERATURE, Kfu

250 300

FIGURE 5 - Ultimate Flexural Strength, a , of Graphi te/Epoxy Composites

203
xlO

300

250

200

150

100

50

{!)

(1) Boron/Epoxy (SP-272)(Ref .58)

(2) HT-S/Polyimide (Ref. 11,58)

(3) HT-S/Epoxy-Phenolic (Ref.

(4) GY-70/Polymide (Ref. 11)

Minimal Data

xlO

20

18

16

8

14

12

10

CM

I

3

50 100 150 200

TEMPERATURE, K fu

250 300

FIGURE 6 - Ultimate Flexural Strength, a , of Miscellaneous Advanced

XlO

50

Composites.

6

40

30

20

10

(1) Thornel 50

(2) HMG-25

z Medium Modulus

Low Modulus

All Ref.

L

(3) Modmor II

UNIAXIAL LONGITUDINAL

(2)__

(3)

J.

10

xlO

30

25 rsi

I E 20 2

15

'^^

10

0 50 100 150 200 250

TEMPERATURE, K

FIGURE 7 - Flexural Modulus, E^, of Advanced Composites.

300

2 04
500

(1) HT-S/X-904 (Ref.l3) (4) Borsic-Ti tanium

^(2) Boron/Epoxy (SP-272, Namco 5505)

Aluminum (6061)^^

(Ref.81) Borsic-Steel/'

Aluminum (6061)

(Ref.l3)

(3) Boron/Aluminum (6061) (Ref . 13)

Minimal Data

(a) UNIAXIAL LONGITUDINAL

60

40 -

20

(b) UNIAXIAL TRANSVERSE

(1) HT-S/X-904 (Ref. 13)

(2) Boron/Epoxy (Narmco 5505,

SP-272) (Ref.81)

(3) Boron/Aluminum (6061) (Ref . 13)

Minimal Data

I

0 50 250

xlO

35

30

25

20

- 15

8

- 10

- 3

- 2

100 150 200 300

TEMPERATURE, K

FIGURE 8 - Ultimate Compressive Strength ,a^", of Advanced Composites.

205
xlO

45

40

35

30

Q.

25

20

15

(a) UNIAXIAL LONGITUDINAL

(1) HT-S/X-904 (Ref.l3)

(2) Boron/Epoxy(Narmco 5505, SP-272) (Ref .81)

(3) Boron/Aluminum (6061) (Ref . 13)

(4) Borsic-Titanium/Aluminum (6061) (Ref . 13)

(5) Borsic-Steel /Aluminum (6061) (Ref .13)

20 |- (1) HT-S/X-904 (Ref. 13)

(2) Boron-Epoxy (Narmco 5505, SP-272) (Ref .81)

(3) Boron/Aluminum (6061) (Ref . 13)

10 h ini IIIVIUXIUI I kfUIMN V (- KNh

(2)

(4)

(1)

(b) UNIAXIAL TRANSVERSE

HI

50 100 150 200

TEMPERATURE, K

250

-•

xlO

30

28

26

24

22

20

18

16

14

12

10

H 16

12

8

-I 4

0

300

FIGURE 9 - Compressive Modulus, E , of Advanced Composites,

206
8

7

6

5

4

3

2

1

0

6

5

3 -

2 -

(a)

HT-S/X-904

at 77K

Avge,

I I I I I J_JL

11 13 15 17 19

SI .SI a , psi xlO-

(b)

HT-S/X-904

Aa

295K^77K

I I I I \ L n

-7.5 -3.5 3.5

A SI

Ao , psi

7.5 . xlO

FIGURE 10 - Histograms illustrating the range of longitudinal

interlaminar shear strength, o^^ , as reported for HT-S/X-904

Graphi te-Epoxy at 77K (a) and the reported changes in

interlaminar shear strength in the same composite upon

cooling from 295K to 77K (b).

207
xlO-"

22

20

18

16

14

^ 12

10

4 -

- (1) GY-70 (Ref.l3)

^(2) Thornel 50 (Ref.8)

\(3) HMG-50 (Ref.lO) \ (4) HM-S (Ref.B)

_

(5) Modmor I (Ref.B)

(6) Fibralloy 300 {Ref.l3)

(7) Samco 320 (Ref.B)

(B) HT-S (Ref.lO, 13,58)

(9) Modmor II (Ref.l3)

(10) HMG-25 (Ref.B)

(11) Thornel 25 (Ref.47)

— High Modulus

Medium Modulus — Low Modulus

Minimal Data

(8) \

(9)

'

(11)'

- 12

xlO'

14

- 10

f E

B

'

150 200

TEMPERATURE, K

250 300

FIGURE 11 - Interlaminar Shear Strength, a^^, Graphi te-Epoxy Composites

Data Reflects Highest Values Reported (See text for details). Flat

Short-beam Test Data except as noted.

2 08
20 F

18

16

14

12 -

10 -

8 -

6 -

4 -

(1) PRD-49/Epoxy (Ref.59,60)

(2) Boron/Epoxy (SP-272) (Ref. 13,58)

(3) HT-S/Polymide (Ref. 58)

(4) Boron/Aluminum (6061) (Ref . 13)

(5) HT-S/Epoxy-Phenolic (Ref. 58)

(6) Borsic-Titanium/Aluminum (6061) (Ref . 13)

(7) Borsic-Steel/Aluminum (6061) (Ref . 13)

(1) NOL-SB

50 100 250

xlO

14

- 12

- 10

- 2

150 200 300

TEMPERATURE, K

FIGURE 12 - Interlaminar Sheer Strength, a^^, of Miscellaneous Advanced

Composites. Flat Short-beam Test Data except where noted.

209
xlO

20

18

16

14

12

2 10

CjJ

-3

(1) HT-S/X-904 (Ref. 13)

(2) 4.2 mil Borsic/Aluminum (6061) (Ref. 13)

(3) 5.6 mil Boron/Aluminum (6061)(Ref. 13)

(4) 4.2 mil Borsic-Titanium/Aluminum (6061) (Ref. 13)

(5) 4.2 mil Borsic-Steel/Aluminum (6061)(Ref. 13)

Longitudinal

Transverse

(5)

UNIAXIAL TENSILE

FIGURE 13

50 100 150 200

' . TEMPERATURE, K

Ultimate Tensile Strain.e^", of Advanced Composites

210
211
xlO^

TEMPERATURE, K

FIGURE 15 - Ultimate Bearing Strength, a^'^, of Boron/Aluminum

Composites, Uniaxial Longitudinal Direction.

212
-5

XlO

100

80 -

60

40

20

0

0

-20

-40

-60

-80

-100

(3) Fibralloy 300/X-904(Ref . 13)

(4) HT-S/X-904 (Ref.l3)

(5) PRD49-I/ERLA 4617(Ref .59,60)

(6) PRD49-I/NA^A Resin 2 (Ref. 59,60)

(7) Boron/Epoxy Sp-272

(8) 5.6 mil Boron/Aluminum 6061

(9) 4.2 mil Borsic-Ti /Aluminum

6061

UNIAXIAL

LONGITUDINAL

All Ref. 13

(10) 4.2 mil Borsic-Steel /Aluminum

(6061)

(11) 4.2 mil Borsic/Aluminum(6061)

I

I I L

0 50 100 150 200 250

TEMPERATURE, K FIGURE 16 - Longitudinal Thermal Expansion, AL/L, of Uniaxial

Advanced Composites.

300

213
-160 -

0

13)

- ^

UNIAXIAL

TRANSVERSE

/ (1) Boron / Epoxy (SP-272) (Ref . 13)

(2) Fibralloy 300 / ERLA 4617l/p . / HT- S/ X-904 J^^^^-

/ (3) 4.2 mil Borsic-Ti/Aluminum 6061 (Ref. 13)

/ (4) GY-70 / X-904 (Ref. 13)

(5) Boron / Aluminum 6061 (Ref. 13)

(6) 4.2 mil Borsic-Steel / Aluminum( 6061) (Ref . 13)

(7) PRO 49-I/ERLA 4617(Ref. 59, 60)

(8) PRD49-I/NASA Resin 2 (Ref. 59, 60)

J I I _J \ L

50 100 150 200 250

TEMPERATURE, K FIGURE 17 - Transverse Thermal Expansion, AL/L, of Uniaxial

Advanced Composites.

300

214
15 -

10 -

5 -

0

1.0

0.8

0.6

0.4

0.2

- (1) HTS/X-904 (Ref. 13) / (2)

-

" (2) Thornel 50/Polaris (Ref. 38)

(3) B-Epoxy, Sp-272 (Ref. 13) /

. (4) B-Epoxy, 4.0 Boron/ / —0)

Polaris (Ref. 38) / y/

/ -

UNIAXIAL

LONGITUDINAL

^ " (4)

% •

1 1 1 , (3) ,

•

1

(1) HT-5/X-904 (Ref. 13) - (2) Thornel 50 / Polaris (Ref. 38)

(3) B-Epoxy, Sp=272 (Ref. 13) (2)

UNIAXIAL

TRANSVERSE

1 1 1

^^J^^'ls)

1 0 50 100 150 200 250 300

TEMPERATURE, K

FIGURE 18 - Thermal Conductivities, A, of Advanced Composites

215
xlO-"

12

10

/X-904 (Ref. 13)

. (2) B-Epoxy,(SP-292 Ref. 13)

(3) Boron/6061 Aluminum (Ref. 13) - (4) Borsic-Titanium/Aluminum (Ref. 13)

(5) Borsic- Steel/Aluminum (Ref. 13)

- (6) Graphite/ Phenolic (Ref. 15)

I

cn

- 6

(1)

HT-S

GY-70

50

Aluminum-"^/

100 150 200

TEMPERATURE, K

250 300

FIGURE 19 - Specific Heat, Cp, of Advanced Composites

216
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 1974
Sunmary; Elastic Properties

During the past year, the elastic properties of thirteen alloys have

been studied at cryogenic temperatures. All of the alloys are commercially

available engineering materials that are candidates for low- temperature

applications, Including superconducting machinery. They were studied between

room temperature and liquid-helium temperature (300-4 K) . Previously, very

little low- temperature elastic data existed for most of these materials. Ultrasonic (10 MHz) pulse-echo and pulse-echo-superposltion methods were used to

detemlne the longitudinal and transverse sound-wave velocities, which are re- lated simply via the mass density to the elastic constants. These constants

Include: Young's modulus, shear modulus, bulk modulus (reciprocal compressibility), longitudinal modulus, Polsson's ratio, and elastic Debye temperature.

Measurements were made semi- continuously on cooling and also on heating.

In the previous semi-annual report, data and discussions were given for

four materials: Ti-6A1-4V, Ti-5Al-2.5Sn (see Appendix A), Inconel 600, and

Inconel X-750. In this report, data and discussions (in the form of manuscripts

Intended for publication) are given for eight additional materials: aluminum

alloys 1100, 5083, 7005, and 7075; and austenltic stainless steels AISI 304,

AISI 310, AISI 316, and A286. Also in this report, data are given for invar

(iron 36 nickel). All the stainless steels and invar have low-temperature

magnetic transitions that cause anomalies in the low- temperature elastic properties of these materials. Invar is a particularly complicated material

magnetically and elastically, and it will be discussed in the next report.

The dynamic magneto-elastic behavior of the stainless steels is much simpler

and is Interpreted in terms of the Doring effect.

218
Contents ; Elastic Properties

Page

Manuscript: Low-temperature elastic properties of four

wrought aluminum alloys 220

Manuscript: Low-temperature elastic properties of four

austenltlc stainless steels 243

Elastic properties of Invar (data only) . . . . , 267

Appendix A: Reprint: Elastic properties of two titanium alloys

at low temperatures 273

>

219
LOW-TEMPERATURE ELASTIC PROPERTIES OF FOUR WROUGHT

ALUMINUM ALLOYS

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

Cryogenics Division

' ^/ Institute for Basic Standards

National Bureau of Standards

Boulder, Colorado 80302

Abstract

The elastic properties of four annealer polycrystalline commercial

aluminum alloys were studied between 4 and JOG K using a pulse-superposition method. Results are given for: longitudinal sound velocity,

transverse sound velocity, Young's modulus, shear modulus, bulk modulus

(reciprocal compressibility) , PoissonV ra io, and elastic Debye temperature. The elastic stiffnesses of the alloys increase four to thirteen percent on cooling from room temperature to liquid-helium temperature. The elastic-constant-temperature curves exhibit regular behavior.

Key words: Aluminum; aluminum alloys; bulk modulus; compressibility;

Debye temperature; elastic constants; elasticity; Poisson ratio;

pulse-echo method; sound velocity; Young's modulus.

Contribution of NBS , not subject to copyright.

* NRC-NBS Postdoctoral Research Associate, 1973-4.

t Present address: Dow Chemical USA, Rocky Flats Division, Golden,

Colorado 80401.

220
LOW-TEMPERATURE ELASTIC PROPERTIES

OF FOUR WROUGHT ALUMINUM ALLOYS

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

Cryogenics Division

Institute for Basic Standards

National Bureau of Standards

Boulder, Colorado 80302

1. INTRODUCTION

Aluminum alloys are used extensively at cryogenic temperatures

because of their favorable mechanical properties. These properties

include: increased strength without loss of ductility at lower temperatures; absence of a ductile-brittle fracture transition; and, for some alloys, high strength-to-weight ratios.

Knowledge of a material's elastic constants is essential to understanding its mechanical behavior. Most mechanical behavior is best

described by a dislocation model, and the elastic constants (usually the

shear modulus and Poisson's ratio) occur in most equations describing

the stress-strain state of a dislocated solid.

In this paper, the elastic properties of four wrought aluminum

alloys — commonly designated 1100, 5083, 7005, and 7075 — are reported

over the temperature range 300 to 4 K. These properties include: the

longitudinal modulus. Young's modulus, the shear modulus, the bulk

modulus (reciprocal compressibility), and Poisson's ratio. While

the changes of the elastic constants in this temperature range are

only moderate, four to thirteen percent, exact values of the elastic

constants are very useful design parameters, permitting accurate

calculations of deflections for any combination of stress and temperature.

Low-temperature elastic constants are also quite valuable theoretically;

they permit the calculation of the Debye characteristic temperature, which

is related in turn to a wide variety of solid-state phenomena that depends on the vibrational properties of solids.

221
An ultrasonic (10 MHz) pulse-superposition method was used for

determining the velocity v of a sound pulse propagated through the

2 specimen. The elastic modulus C is then given by C = pv , where p is the mass density. Different elastic constants were determined from

different modes of ultrasonic excitation. This method has many advantages: small specimens are sufficient; thus, ancillary equipment

such as probes and dewars can also be small, and refrigeration costs

are low; specimens can have a simple geometry and can be easily prepared; measurements can be made as nearly continuously as desired;

relative precision is high, about one part in 10^ for the velocities;

laboratory-to-laboratory variations of the elastic constants are typically

a few percent or less; and tests are completely non-destructive.

Low-temperature elastic data for aluminum alloys have two-fold

interest. First, the elastic constants provide basic information about

interatomic forces. Second, the same numbers are essential design parameters for stress-bearing members. Data given here permit the loaddeflection behavior of the alloys to be accurately predicted between

room temperature and liquid-helium temperature. Accurate elastic data

become especially important in applications involving high stresses,

large structural parts, or precision parts.

222
2. MATERIALS

Aluminum alloy 1100 is commercial-quality aluminum. It has good

corrosion resistance, high electrical and thermal conductivities, high

ductility, but low strength properties. Strength can be improved somewhat

by strain hardening without significantly decreasing other properties.

The main impurities in this alloy are usually iron and silicon.

Aluminum alloy 5083 is characterized by good welding properties

and by good corrosion resistance in marine environment. Magnesium is the

major alloying element and, along with manganese, produces a moderately

strong, yet ductile, alloy, which does not respond to heat treatment.

The strength properties of 5083 improve with lower temperatures. The

main advantage of 5083 seems to be its weldability; welds as strong as the base metal can be obtained. Currently this alloy is being used in

a number of applications involving the manufacture, transfer, and storage

of liquefied natural gas; these applications require many millions of

kilograms of materials.

223
Aluminum alloy 7005 is a heat-treatable alloy containing zinc and

magnesium, which are balanced to obtain a natural- aging alloy. Chromium is added to reduce corrosion of the heat-affected weld zones, and

zirconium is added to reduce weld cracking and to improve mechanical

properties.

Aluminum alloy 7075 contains zinc as the major alloying element.

Together with a small percentage of magnesium, this alloy can be precipitation hardened to produce high strength.

Details of compositions, heat treatments, mass densities, and hard- nesses of the alloys are given in Tables I and II.

224
3. EXPERIMENTAL

Alloys were obtained from commercial sources; 1100 and 7075 in

the form of 3/A-in. (1.9-cm) rods, 5083 and 7005 in the form of 3/4-in.

(1.9-cm) thick plate. Cylindrical specimens 5/8-in. (1.6-cm) in diameter and 5/8-in. (1.6-cm) long were prepared by grinding. Opposite

-4

faces were flat and parallel within 10 in. (2.5 ^m) . Specimens were

annealed at a pressure of 5 x 10 ^ torr, or less, and cooled in the

furnace. Hardnesses were determined by standard metallurgical methods,

and mass densities were determined by Archimedes 's method using distilled

water as a standard.

Quartz transducers (10 MHz) were bonded to the specimens with

phenyl salicylate for room- temperature measurements and with a stopcock

grease for lower temperatures. In a few cases, failure of these bonds

at very low teii:5)eratures required using a silicone fluid (viscosity = 200,000 cP at 25 °C) for bonding. The low- temperature apparatus was described previously fl].

A pulse-superposition method was used to determine the sound-wave

velocities over the temperature range 300-4 K. Details concerning this

method were given elsewhere [2].

225
4. RESULTS

Quantities that were measured directly are the longitudinal and

the transverse soimd-wave velocities Vg and v^. From these, the

longitudinal modulus Cg and the transverse modulus were calculated

according to

and

^ ^t = P^t •

These moduli are shown in Figs. 1 and 2 for the temperature range studied.

Errors in the absolute velocities are believed to be about one-half percent or less. All of the other elastic constants that are used to describe polycrystalline aggregates are simply related to these two moduli.

The moduli considered here — the shear modulus G, Young's modulus E, the

bulk modulus B, and Poisson's ratio v are given by:

G = C^, (3)

E = 3G(Cg -

-| Cj.)/(Cg - C^.) , (4)

B = Cg - I

C^. (5)

and

V = I

(Cg - 2C^)/(Cg - C^). (6)

The elastic constants obtained from these relationships are shown as functions of temperature in Figs. 3-5. Values of the elastic constants

at selected temperatures are given in Table III.

The temperature variations of the elastic constants can be described

mathematically in various ways. In this case, the temperature dependences

of both Cg and C^. were fitted to a theoretical relationship suggested

by Varshni [3] :

C = C»-^7|r— , (7) e - 1

226
where C is any elastic constant (Cg and in this case), C°, s, and t 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 IV. Room-temperature values of the temperature coefficients

of the elastic moduli are given in Table V; these values occur in the

linear high- temperature region. Average differences between measured

and curve values are 0.03% and 0.06% for the longitudinal and transverse

moduli, respectively.

The elastic Debye temperature 6 can be calculated from the elastic wave velocities by [5]

e = K < V >, (8)

where

K = ll A) (9)

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

-3 -3

vg + 2v

< V > = — . (10)

The elastic Debye temperatures for the four alloys at T = 0 K, and

also for unalloyed aluminum are given in Table VI.

227
5. DISCUSSION

As shown by the data in Table VI, changes in the elastic constants

of aluminum and its alloys between 300 and A K are about four percent

for B, twelve percent for E, twelve percent for G, and four percent

for V. These changes are somewhat larger than those observed in alloys

based on copper or iron, for example. Most of the changes occur above

about 100 K. Below this temperature the elastic constants change only

slightly with tenqjerature. Thus, any changes in the mechanical behavior

of these alloys in this temperature region probably cannot be ascribed

to an elastic origin.

The temperature behavior of the elastic constants of the alumlnxam

alloys reported on here is quite regular . The regular decrease of the

moduli with increasing temperature and the relative flatness of the

modulus- temperature curves at lower temperatures has already been mentioned. Also, in accord with the third law of thermodynamics, the slopes

dC/dT approach zero at zero temperature. And a linear temperature dependence is exhibited at higher temperatures, those higher than about 150 K,

which is roughly one-third of the Debye temperature. Besides indicating

the absence of magnetic or structural transitions, this ideal temperature

behavior suggests that the anharmonic properties of these alloys can probably be explained by a relatively simple model.

These alloys were not examined for texture. However, texture

would have little effect on their elastic properties. Aluminum single

crystals are only slightly anisotropic; the Zener anisotropy ratio for

aluminum is 1.2; it is 1.0 for the isotropic case. Alximinum's isotropy

is purely accidental since aluminum, because of its three valence electrons,

has a large band-structure contribution to the elastic constants flO].

This contribution is usually anisotropic; the anisotropy is canceled in

the case of aluminum by other energy terms. Thus, since aluminum single

crystals are only slightly anisotropic, even a strongly textured polycrystalline aggregate of aluminum would have nearly isotropic elastic

behavior.

It should be emphasized that the data reported here are dynamic

(adiabatic) rather than static (isothermal); they apply strictly to

rapid rather than slow loading. However, the differences between

228
adiabatic and isothermal elastic constants are small. They become

smaller at lower temperatures because of the diminishing thermalexpansion coefficient, and they vanish at zero temperature. Using

formulas given by Landau and Lifshitz [6] it can be shown for aluminum

at room temperature:

0.045 (11)

0,

where subscripts S and T denote the adiabatic and the isothermal cases,

respectively. For E, B, V, and G, these corrections are typically in

the ratio 1:9:4:0 if v has a value near 1/3.

Effects of alloying on the elastic properties of aluminum cannot be

accurately determined from the present study because of the large number

of alloying elements and their interactions. Such effects have been

considered elsewhere [11]. However, some general observations can be made concerning alloying. Disregarding alloy 1100, with respect to

'pure' aluminum, the shear modulus and Young's modulus increased in all cases while the bulk modulus decreased, and by a larger percentage.

Poisson's ratio decreased in all four alloys. The 1100 alloy, which

contains only one percent of Impurities, presents an interesting case. The data indicate that while the shear modulus of this alloy is identical

to that of unalloyed aluminum, the bulk modulus is higher by about five

percent. Since the bulk modulus is not measured directly, but is calculated

from the difference of two velocities according to Equation 5, a compounding

of errors may account for this discrepancy. If the effect is real, then

it has important consequences for the problem of averaging single-crystal

elastic coefficients to obtain the bulk modulus of a polycrystalline

aggregate.

Finally, approximate relationships among the elastic constants

are indicated. For all of the alloys, and for all temperatures, as a first approximation:

= 0.005

S T

= 0.020 , and

229
and

V ^ 1/3

B S5 E ~ 8/3 G .

(12)

(13)

These should be useful for many engineering purposes where only rough

numbers are needed and only one of the elastic constants is known.

CONCLUSIONS

From the results of this study the following conclusions are drawn:

1. All the elastic properties of aluminum alloys 1100, 5083, 7005,

and 7075 behave regularly with respect to temperature.

2. For all alloys studied, the temperature behavior of both Cg and

C^ can be described accurately by a theoretical relationship suggested

by Varshni.

3. In this series of alloys, alloy 5083 has the highest Young's

modulus, the highest shear (rigidity) modulus, and the lowest Poisson's

ratio. Alloy 7005 has the highest bulk modulus.

ACKNOWLEDGMENT

This work was supported in part by the Advanced Research Projects

Agency of the U.S. Department of Defense.

REFERENCES

1. E. R, Naimon, W. F. Weston, and H. M. Ledbetter, Cryogenics ' 14 (197A) 246.

2. H. J. McSkimin, J. Acoust. Soc. Amer. 33^ (1961) 12.

3. Y. P. Varshni, Phys. Rev. B2 (1970) 3952.

4. G. N. Karam and G. A. Alers, J. Appl. Phys. 35 (1964) 327.

5. P. Debye, Ann. Phys. (Leipz.) 39. (1912) 789.

6. L. D. Landau and E. M. Lifshitz, Theory of Elasticity (Pergamon,

London, 1959) p. 17.

7. T. Lyman (Ed.), Metals Handbook (Am. Soc. for Metals, Metals

Park, Ohio, 1961)

8. R. Develay, A. Faure, S. Lehongre, D. Mugnier, and D. Schroeter,

in Advances in Cryogenic Engineering, Vol. 12, K. Timmerhaus (Ed.)

(Plenum, New York, 1967) p. 484.

9. J. L. Christian and J. F. Watson, in Advances in Cryogenic Engineering,

Vol. 6, K. Timmerhaus (Ed.) (Plenum, New York, 1961) p. 604.

10. T. Suzuki, Phys. Rev. B3 (1971) 4007.

11. W. Koster, Z. Metallk. 32 (1940) 282.

230
List of Tables

I. Compositions of the alloys, mill analyses, wt. pet 232

II. Properties of the alloys 233

III. Values of the elastic constants of aluminum alloys at

11 2 selected temperatures in units of 10 N/m except v,

which is dimensionless 234

IV. Parameters in equation 7 235

V. Temperature coefficients of the elastic constants at room

_A -1

temperature (10 K ) 236

VI. Elastic Debye temperatures at T = 0 K 237

List of Figures

1. Temperature variation of the longitudinal modulus 238

2. Temperature variation of the shear modulus 239

3. Temperature variation of Young's modulus 240

4. Temperature variation of the bulk modulus

(reciprocal compressibility) 241

5. Temperature variation of Poisson's ratio 242

231
o

4J

CO

0)CO

t>^ iH

cdG

CO

COo

iHH

CO

(U u

o

COo

•H

4-1 •H

COo

o

CO H

•HH

O

o

o

CO

O . CM V

C» CO o

COo

o

V

CO

00

3

•

cy\ V

o

O

V

CO m tH CM CO

iH r-l

CO CO CO CO PQ PQ CQ pq

O CO m in O oo o r-i O o o iH 232
Table II. Properties of the alloys

Alloy Hardness

(DPH No .

,

1 kg load)

Mass density ^ at 294 K, g/cmCondition

1100 28 2.818 Annealed 345**C; furnace cooled

5083 78 2.666 Annealed 413°C, 1/2 h; furnace cooled

7005 77 2.779 Annealed 400''C, 3 h; furnace cooled

7075 67 2.721 Annealed 413**C, 3 h; furnace cooled

233
> o

4J

(0

a

(Ua

10g

a

CI 2 o

•H iH U i-l CO O

n) r-( iH

(U M-lO

0)

CO

4-1

U

•H

M-l0

to

OJ3

tH

(0 >

fn > o

00

O in

CN

<|—

CO

J-l iH3

<n

(U

CO3 o\ 0 00

•H VD >

(U

a>

(Q

(U

l-i 3

n) ^ o h om aH

<uo

4)

l-i CU

l4-{ (U

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CD o in

4J iH3

COV

oiu

sen

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w

vO

-<r CQ <N

(11 M3

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i"

OJ H0 o i-H rH <

fo

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n CO

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en 00 00

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o\ o

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cnoocOiH r^csir^vo

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in <3V CM VO

\© On o\

CM CM CM CM o in 00 CO vo CO

CM CN CM <M

r~- r- O O O O

inoo^ooN incor^ON iHinoNO o <! r-- CO r~r-~r-oo r~r^r-«r^

CM 00 CO

O ~3- r-~ 00

r-^

iH 00 1-1 -3- O CO r-~ 00

r-» r~-

\D cr» ON CM

tH CM CO -3-

r~. 1-^ o o o o o o O CO CM iH

CO

00om

ON in o < vo r~ 00 r-oooo<-( iHP-O^ OiHCMCO ^Or^ONON o o o o o o o CO CM iH O O O O O

CO <N r-l

OO

O

O

CO CM rH *e moo n

ulao

<4H o

CU (0 ua e

<u CO 4J

OJ CO

M-l<U 60 UG

•H 3

CO 4J

CO >

T3 •HM

rH o;

CO T3 U

<n 0) >^ M u 0) u 3 1<u CO rH OJ 00 CO c 01 •H J3

CO 4Jc a OJ 0 I-l u CO

14-1

a

T3 c

01 •H

4J

CO CO rH OJ 3 •H O I-l rH 4-1 CO cM

HCO

CM
Table IV. Parameters in equation 7

Alloy Mode s t

10-'--4^/m^ K

1100 £ 1.136 0.0903 255.9

t 0.2892 0.0301 206.9

5083 9. 1.151 0.0926 235.3

t 0.3067 0.0381 206.7

7005 1.175 0.0928 240.0

t 0.2963 0.0327 203.4

7075 S. 1. 126 0.0707 219.6

t 0.2961 0.0344 223.3

235
Table V. Temperature coefficients of the elastic constants

-4 -1

at room temperature (10 K )

All Alloy

1 dB 1 dE 1 dG 1 dv

B dT E dT G dT V dT

1100 -2.01 -5.01 -5.37 1.41

5083 -1.92 -6.06 -6.58 2.08

7005 -2.13 -5.48 -5.88 1.54

7075 -1.58 -5.08 -5.52 1.75

Aluminum* -1.97 -5.53 -5.77 1.80

* Derived from single-crystal data in reference [4].

236
Table VI. Elastic Debye temperatures at T = 0 K

Alloy 6(K)

1100 426.2

5083 440.4

7005 425.9

7075 422.2

Aluminum 430.6*

* Calculated from single-crystal data

in reference [4].

237
00C\J<£)O^00CSJ<^O«a-00 I—I 1—ii—ii-H If—ii—I oooo

(^ui/N^^oT) smnaow ivNianii9N0i

238
on

cc. Ul

a.

UJ

00 o CM 00 o CM CM 1—

«

o o cn If) ro ro CM CM CM CM CM CM CM o O o o C3 o o o C3 snmaow yv3HS

239
240
m

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a

ou

1—1

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0)

^1

cn

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0)

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-i->

0)

(^ui/N^^OT) snmaow >nn9

241
1 r—r—I—1

1—I—1

:t

o U3 CM 00 O U3 CM 00 o ID CO CO CM CM r-l t—

r-l CO ro CO m CO m cn CO CO O o O O o o o CD Q o CD oim s.Nossiod

242
LOW-TEMPERATURE ELASTIC PROPERTIES OF FOUR AUSTENITIC STAINLESS STEELS*

H. M. Ledbetter, W. F. Weston*, and E. R. Nalmon*t

Cryogenics Division, Institute for Basic Standards,

National Bureau of Standards, Boulder, Colorado 80302

Abstract

The elastic properties of four austenitic stainless steels — AISI 304,

AISI 310, AISI 316, and A286 — are reported over the temperature range 300 to 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 dilatational component,

decrease anomalously at tetperatures below 80 K. The largest anomaly, about 3%,

is in the bulk modulus of the 304 alloy; this modulus is lower at 0 K than at

300 K. Results are interpreted on the basis of the Dbring 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.

Key words: Bulk modulus; chromium alloys; compressibility; Debye temperature;

Dbring effect; elastic constants; iron alloys; magnetic transition; nickel alloys;

Poisson's ratio; pulse-echo method; sound velocity; Young's modulus.

*

t

Contribution of NBS, not subject to copyright.

NRC-NBS Postdoctoral Research Associate, 1973-4.

Present address: Dow Chemical USA, Rocky Flats Div. , Golden, CO 80401.

243
LOW-TEMPERATURE ELASTIC PROPERTIES OF FOUR AUSTENITIC STAINLESS STEELS

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

Cryogenics Division, Institute for Basic Standards, National Bureau of

Standards, Boulder, Colorado 80302

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 become 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 low temperatures. For example, lower temperatures may promote a change

of the crystal structure from fee to body-centered cubic (bee) or to close-packed

hexagonal (eph) , thus probably embrittling the material. Other changes such as atomic ordering or magnetic ordering may also occur at low temperatures; these

also affect mechanical behavior. Thus, the nature and the magnitude of a material's low-temperature properties cannot be predicted a priori from room- temperature observations, and there is no substitute for careful low- temperature

experimental determinations of the important properties of each material of

interest.

Oilman''" concluded that "the most important mechanical characteristic of

a crystal is its elastic modulus." While its significance may often be obscured

by impurities, mierostructure, grain boundaries, residual stresses and strains,

and temperature, it nevertheless remains fundamental to a material's mechanical

behavior. Trends in elastic constants can be correlated with phase instabilities;

changes in crystal structure are generally accompanied by discontinuities in the

elastic constants. Higher-order phase transitions such as order-disorder and

nonmagnetic-magnetic also exhibit elastic anomalies. From elastic constant data,

information on interatomic potentials can be deduced; similarly, the validity of

proposed interatomic potentials can be accurately tested by calculating elastic

constants from them and comparing results with observation. By invoking a model.

244
interatomic force constants can be derived from the elastic constants; the

force constants are fundamental in that they describe directly the strength

of interatomic "bonds." Debye characteristic temperatures can be calculated

simply from elastic constant data; the Debye temperature relates in turn to a wide variety of solid-state phenomena including atomic diffusion, lattice

thermal conductivity, sound velocities, vibrational entropies, zero-point

energies, electrical conductivity, Bragg-ref lection intensities, atomic vibrational amplitudes, and melting. Also, Gruneisen parameters can be calculated

from higher-order elastic constants; these relate in turn to a variety of

anharmonic solid-state properties, such as thermal expansion and attenuation

of acoustic waves in solids. Besides the examples cited above, other metallurgical

problem areas where elastic constants play a key role include: theories of solidsolution strengthening, alloy development, models of maximum ultimate strength,

and dislocation models of plastic behavior, a dislocation description often involves the shear modulus and Poisson's ratio.

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* — 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 antif erromagnetic states. A magneto-elastic interpretation of the

anomalies is given.

* This tradename is used to describe the material; its use does not imply an endorsement by NBS of a particular product.

245
II. EXPERIMENTAL

A. Specimens

Materials were obtained from commercial sources In the form of 3/4 - In

(1.9 - cm) diameter rods. Their chemical conpositlons 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

-4

cylinders 1/2 - in (1.2 - cm) thick with faces flat and parallel within 10 in

(2.5 ^im).

B. Procedures

2 3 A pulse-superposition ' method was used to measure the longitudinal and

transverse sound-wave velocities between room temperature and liquid-helium

4

temperature. The specimen holder, which was described previously , was placed

in the ullage of a helium dewar and lowered (raised) stepwise to achieve cooling

(heating). Measurements were made semi- continuously on cooling, and a few points were checked on heating to verify reversibility. Temperatures were measured by

a chromel-constantan thermocouple contacting the specimen. 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 thermal- 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 estimated to be about one percent. The imprecision in the relative velocities is a few parts in 10^.

III. RESULTS

Longitudinal and transverse moduli are shown in Figs. 1 and 2 as a function

of temperature. The longitudinal modulus Cg is given by

= pvg . (1)

where p is the mass density and Vg is the longitudinal wave velocity. The transverse

modulus C is identically eqioal to the shear modulus G and is given by

246
= G = pv^ , (2)

where is the transverse wave velocity. Young's modulus E is given by

and is shown in Fig. 3. The bulk modulus B, or reciprocal compressibility, is

given by

B = - I

(4)

and is shown in Fig. 4. Poisson's ratio v is given by

V =1 (C^ - 2C^)/(C^ - (5)

and is shown in Fig. 5.

Temperature dependences of both C„ and C were fitted to a semi- theoretical

5

^ relationship suggested by Varshni :

C = C° - (6)

e^' - 1

where C°, s, and t 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 III. Of course,

the low- temperature elastic anomalies are not described by Eq. (6) , which describes

the extrapolated higher- temperature behavior shown as dashed lines in the figures.

Temperature coefficients of the elastic constants at room- temperature are given

in Table IV.

Elastic Debye temperatures were calculated 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)

247
where

Here h is Planck's constant, k is Boltzmann' s constant, N is Avogadro's number,

p is the mass density, and A is the effective atomic weight. The average velocity

is given by

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 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^ Resonance tests on AISI 303 for E and G simultaneously

at temperatures between 320 and 4 K showed anomalies below 80 All these anomalous

results can be interpreted in terms of the usual AE = AE. effect (discussed below)

A, 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. Stainlesssteel-type alloys have been shoxro to be antiferromagnetic at lower temperatures;

' 6 8 Neel temperatures near 40 K have been reported ' 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 difference in the 304 alloy between the Neel temperature (40 K) and the

Z48
temperature below which anomalous elastic behavior is observed (80 K) . Short9

range magnetic-order effects have been discussed by Schlosser for face-centered

cubic 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, showing 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 secondorder type.

The usual so-called AE effect, better designated due to linear magnetostriction strains induced by an applied stress, and usually interpreted by invoking

a Weiss domain model (with domain rotations and domain-wall 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?""'" Also, effects due to AE. are not

12 observed at high frequencies 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, B, and v without affecting G is the effect

due to spontaneous voliime magnetostriction. This effect is designated AE^, where

CO = (V - ^p)/^p is the spontaneous volume magnetostriction, V is the total volume,

and V is the volume in the paramagnetic state. The effect was first reported by

^13 .. 14

Engler for an Fe-42 Ni alloy. It was explained first by Doring 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 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:

249
AE^E^ = (1/E)^ - (1/E)jj = - I

Oa)/aH)^^^/OM/8H)^^^ , (10)

where BoVdH Is the forced volume magnetostriction due to the magnetization,

2 9M/BH is the high-field susceptibility x> E is E^^, and a is the tensile stress.

An alternative derivation of Eq. (10) was given by Hausch^"^ The factor of 1/9

2 16 in Eq. (10) is really (1 - 2v) where v is Poisson's ratio ; thus, 1/9 applies

strictly only when v = 1/3. From the data in Fig. 5, v ranges from 0.27 to 0.31

for the materials and temperatures of interest here; thus 1/9 should be replaced

by 0.17 as an average value for present purposes. It should be noted that the

Doring effect always lowers the moduli (E^^ ^ » whether the volume magnetostriction is positive or negative. This is related to the fact that a relaxation

mechanism is Involved in going to an antiferromagnetic state, and such mechanisms

always soften the elastic stiffnesses. The Doring effect in iron-nickel alloys

17 was discussed by Koster, who denoted it as AE. rather than as AE . The effect

has been invoked by several authors to explain elastic anomalies in magnetic materials,

usually of the invar type. Hausch'''^ criticized the relevance of the Doring effect

18 in most of these cases. Schlosser discussed the magneto-volume contribution to the

compressibility of invar, but Hausch'^^ 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 AE^ 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}^ an iron alloy

containing about 35 pet. Ni and having the same crystal structure as the alloys

2 _ 2 considered here, it is calculated from Eq. (10) that ^^^E = - 1.4 x 10 cm /dyn.

For the 304 stainless steel alloy, the present results give for the bulk modulus

2 — lA 2 AE/E = - 0.2 X 10 cm /dyn. 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 in invar-type

alloys. Besides invar and other iron-nickel alloys, large magneto-volume effects

250
have also been observed in Iron-manganese, iron-palladium, iron-platinum, and

17

19 iron-cobalt alloys. A quantitative correspondence is also precluded because

the alloys reported on here were studied in mechanically deformed states. Koster

showed that the AE effect depends sensitively on metallurgical variables. The

usual AE = AEj^ effect is usually suppressed by mechanical deformation. No studies seem to have been made on the effect of mechanical deformation on AE . It would

CO be expected that the AE effect will be enhanced in annealed materials since the

residual stresses due to mechanical deformation would not interfere with domainwall motions.

The best parameter for evaluating the Dbring effect is the bulk modulus,

which characterizes purely dilatational elastic distortions. Shear distortions

are isovolumic and should not change the magnitude of the magnetization. Young's

modulus and Poisson's ratio are less satisfactory parameters in this case because

they depend more on shear distortions than on dilatational distortions. Using

the familiar relationship

E = 9GB/(G + 3B), (11)

it follows from Eq. (10) that the defect in the bulk modulus due to the Doring

effect is

ABJB^ = (1/B)j^ - (1/B)jj = - 9(1 - 2v)2(9co/aH)^^^/(aM/aH)^^^, (12)

2 where B = B B and P is pressure. And from the relationship

rl M

V = fE/(2G)] - 1, (13)

it follows that

2 Av^v^ = - ^+^^ (aavaH)^^^/(aM/aH)^^^, (i4)

V E

and, of course, AG^20 =0.

Alers, Neighbors, and Sato have discussed the Doring effect as arising

from the change in spontaneous magnetization due to volume changes resulting

from the measurement process itself. Thus, this effect is not an intrinsic part

251
of the magnetization change, and it is not related f tmdamentally to the interatomic magnetic interaction. This is also clear from the work of Hausch}^ who

considered the Dbring effect to be due to a stress-induced change in the magnetization intensity. (The usual fundamental magnetic energy terms are the exchange,

21 anisotropy, magnetoelastic , and magnetos tatic ; the latter vanishes for zero external field.) Since the elastic constants only have a practical meaning when

they are being "measured," either in the laboratory or in a stress-bearing structure,

the viewpoint of Alers et al. does not mean that the dashed curves in the figures are a more appropriate measure of the elastic constants. The thrust of the con- siderations of Alers et al. is that there is no fundamental interatomic interaction that would, if accounted for, predict the Dbring effect. Recall that any

energy term, magnetic or not, contributes to the elastic constants through its

second spatial derivative. For the Doring effect there is an "extra" or "extrinsic"

negative elastic constant that is not related to a fundamental energy term of

the material. The more familiar AE = effect is also not a fundamental magnetic

effect in the sense discussed here.

which is a standard thermodynamic result, where a = (1/V) (dW/dT) and C = TOS/BT) . P P P

This emphasizes the value of the thermodynamic approach to phenomena such as the

Dbring effect and suggests other thermodjmamic correlations. For example, from

the well-known relationship

(15)

(16)

where C = T(9S/9T) , the corresponding relationship in magnetic variables is

(17)

Thus, a model connecting x and the constant-pressure and constant-volume

susceptibilities, would permit an estimate of the magnitude of AB .

252
It was suggested by Hayes and Miodownlk that both AE^ and AB^ are proportional to Bj^. The present data do not support this suggestion, but the effect

could be masked by random errors arising from inhomogeneities, mechanical deformation, absence of a technically saturating magnetic field, etc.

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

antiferromagnetic phase.

The desirability of further studies on these alloys, especially with respect

to the Coring effect, is indicated. 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

24 damping mechanism in ferromagnetic materials.

V. CONCLUSIONS

From the results of the present study the following conclusions are drawn:

1. Stainless-steel alloys AISI 304, AISI 310, AISI 316, and A286 have qualitatively

similar elastic-property variations with temperature.

2. The shear modulus behaves regularly over the entire temperature range studied.

3. The elastic constants with dilatational components — Young's modulus, the

longitudinal modulus, the bulk modulus, and Poisson's ratio — behave

regularly above about 80 K, but anomalously at lower temperatures. The

anomalies are largest for the bulk modulus and smallest for Young's modulus.

4. 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.

5. As suggested first by Doring, the anomalies can be interpreted thermo dynamically

as the difference between constant-magnetization and constant-field bulk moduli.

253
ACKNOWLEDGMENT

This work was supported in part by the Advanced Research Projects Agency

of the Department of Defense.

REFERENCES

1. J. J. Oilman, Austral. J. Phys. l^, 327 (1960).

2. H. J. McSkimin, J. Acoust. Soc. Amer. 33., 12 (1961).

3. H. J. McSkimin and P. Andreatch, J. Acoust. Soc. Amer. 3^, 609 (1962).

4. E. R. Naimon, W. F. Weston, and H. M. Ledbetter, Cryogenics lA, 246 (1974).

5. Y. P. Varshni, Phys. Rev. B2, 3952 (1970).

6. R. P. Mikesell and R. P. Reed, J. Res. Nat. Bur. Stds. (U.S.) 70C , 207 (1966).

7. P. E. Armstrong and D. T. Eash, in Advances in Cryogenic Engineering, Vol. 14

(K. D. Timmerhaus, ed.. Plenum, New York, 1959), p. 64.

8. E. I. Kondorsky and V. L. Sedov, Sov. Phys. JETP 35 , 1104 (1959).

9. W. F. Schlosser, J. Phys. Chem. Solids 32, 939 (1971).

10. J. W. Christian, The Theory of Transformations in Metals and Alloys

(Pergamon, Oxford, 1965), p. 214.

11. W. F. Brown, Jr., Phys. Rev. 50, 1165 (1936).

12. R. M. Bozorth, W. P. Mason, and H. J. McSkimin, Bell Syst. Tech. Jour. 30,

970 (1951)

13. 0. Engler, Ann. Phys. Leipz. 31, 145 (1938).

14. W. Dbring, Ann. Phys. Leipz. 32, 465 (1938).

15. G. Hausch, Phys. Status solidi (a) 15, 501 (1973).

16. E. P. Wohlfarth, Phys. Status Solidi (a) 10, K39 (1972).

17. W. Koster, Z. Metallkde. 35, 194 (1943).

18. W. F. Schlosser, Phys. Status Solidi (a) 18, 235 (1973).

19. W, F. Schlosser, Intern. J. Magnetism 2^, 167 (1972).

254
20. G. A. Alers, J. R. Neighbours, and H. Sato, J. Phys. Chem. Solids 13,

40 (1960).

21. C. Kittel, Rev. Mod. Phys. 21, 541 (1949).

22. L. D. Landau and E. M. Lifshitz, Theory of Elasticity (Pergamon, London,

1959), p. 17.

23. E. J. Hayes and A. P. Miodownik, Phys. Status Solidi (a) 10, K43 (1972).

24. A. Cochardt, Trans. Amer. Soc. Mech. Engnrs. ^l* ^196 (1953).

255
LIST OF TABLES

I. Chemical analyses of the alloys, wt. pet., obtained from mill

analyses 257

II. Densities and hardnesses of the alloys 258

III. Parameters in the temperature-dependence equation (6) 259

IV. Temperature coefficients of the elastic constants at room temperature

(10"^ k""*-) 260

V. Elastic Debye temperatures at T = 0 K 261

LIST OF FIGURES

1. Temperature dependence of the longitudinal moduli of four stainlesssteel alloys 262

2. Temperature dependence of the transverse (shear) moduli of four stainlesssteel alloys 263

3. Temperature dependence of the Young's moduli of four stainless-steel

alloys 264

4. Temperature dependence of the bulk moduli (reciprocal compressibilities)

of four stainless-steel alloys 265

5. Temperature dependence of Poisson's ratios of four stainless-steel

alloys 266

256
o

1—1 r—

(

<;

1—1 1—1 1—

)

r-tnJ

M-l mM-l mw mw

>

ro

I—

<

•rH •

•r-l • • • •

1—

1

in o O O o

v" ) o o o o

•

CO

• • o•

1—

1

• un

r—

1

00

1

• « • • I—

1

t—

r

1—

1

1—

1

o

1—

1

1—1 00

:^

• (M

•

1—

1

1—

1

U •

• 00• 00• 00

00 vO

00 IT) o o O o

00

O

o 1—

1

1—

1

ro CO CO

00

00
Table 11. Densities and hardnesses of the alloys

Alloy Mass density Hardness

(g/cm^) (DPHN, 1 kg load)

304 7.86 225

310 7. 85 220

316 7.97 210

A286 7.95 , 275

258
Table III, Parameters in the t emperature- dependence equation (6)

Alloy Mode

(lO^^N/m^)

s (lO^^N/m^)

t

(K)

304 i 2. 608 0.297 371.8

t 0.803 0.089 251. 8

310 a 2.647 0. 175 271. 8

t 0. 767 0. 051 173. 7

316 3.028 0. 332 369. 7

t 0.914 0.099 243. 5

A286 a 2.636 0. 334 454.4

t 0. 809 0.080 243. 1

259
Table IV, Temperature coefficients of the elastic constants at room

-4 -1

temperature (10 K )

Alloy 1 dB 1 dE 1 dG J_ dv

B dT E dT G dT V dT

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

260
Table V. Elastic Debye temperatures at T = 0 K

Alloy e (K) e (K)

magnetic extrapolated paramagnetic

304 467.8 468.1

310 - 459.0 459.0

316 496.7 496.8

A286 467.6 467.7

Iron (bcc)^ 472.4

Chromium^ 453,0

Nickel^ 476.0

^ R. Wanner, Can. J. Phys. 48, 1270 (1970).

F. H. Herbstein, Adv. Phys. J^, 313 (1961).

261
o

r

00 C\J UD o 00 o n LD CO O un o CM Ln

1—

o CD cr> CO un

CO CXI oo CM CM CM CM CM CM (^lij/M^^OT) smnaow ivNianiisNoi

262
00 CO CO CO CO CO 00 00 CO CO CO

Lf) CO o 1—

1

CO LD 00 CO CO CO r-. X)

o O o o o o o o o O o

(^m/N^^oi) snmaow mns

263
T

^1

o CO CM O CO CM o Ln o CO cn .—1 O cn

cn m CM 1—

o o CO CO

CNJ CM CM CM CM CM r—t 1—1 .—

1—

(^uj/N^^OT) snmaow s-snooa

264
265
266
ELASTIC PROPERTIES OF INVAR

List of Figures — Invar

Page

1. Temperature dependence of the longitudinal modulus of

invar (Fe-35Ni) 268

2. Temperature dependence of the shear modulus of invar (Fe-35Ni) . . . 269

3. Temperature dependence of Young's modulus of invar (Fe-35Ni) .... 270

4. Temperature dependence of the bulk modulus of invar (Fe-35Ni) .... 271

5. Temperature dependence of Poisson's ratio of invar (Fe-35Ni) .... 272

267
2.060 1 : 1 —

0 50 100 150 200 250

TEMPERATURE (K)

Figure 1. Temperature dependence of the longitudinal modulus of

invar (Fe-35Ni)

268
0 50 100 150 200 250 300

TEMPERATURE (K)

Figure 2. Temperature dependence of the shear modulus of invar (Fe-35Ni)

269
270
Z71
272
i

i

Appendix A

I
Sound velocities and elastic constants were determined semi-continuously for two

annealed polycrystalline titanium alloys between 4 and 300 K. Results are given for: longitudinal sound velocity, transverse sound velocity. Young's modulus, shear modulus,

bulk modulus, Poisson's ratio, and elastic Debye temperature. A pulse-superposition

technique was used.

Elastic properties of two titanium alloys at low

temperatures

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

Despite the fact that titanium alloys Ti-6A1-4V and

Ti-5Al-2.5Sn are intended mainly for high strength, high

temperature applications, they also have cryogenic uses. Their important properties include: ease of fabrication, corrosion resistance, high strength-to-density ratios (especially at cryogenic temperatures), and high strength-tothermal-conductivity ratios. , . . The low temperature elastic properties of these alloys are reported here. A pulse-superposition method was used to determine the ultrasonic wave velocities in annealed

specimens. From longitudinal and transverse wave velocities, together with the mass density, the elastic constants were

calculated: Young's modulus, shear modulus, bulk modulus

(reciprocal compressibility), Poisson's ratio, and the elastic Debye temperature. These constants are useful in a wide

variety of applications, from engineering design to equations

of state for solids.

A pulse-superposition method was used to determine the sound-wave velocities over the temperature range 4-300 K. No thermal contraction corrections were made; for titanium

alloys this introduces a maximum error (over a 300 K range) of about 0.2??. No bond corrections were made; this error

is insignificant for the purposes of the present study. Maximum uncertainties in the velocity measurements are estimated

to be about 1%.

Results

Longitudinal and transverse sound-wave velocity data are shown in Figs 2 and 3, where the longitudinal modulus is given by

4

C\ = pvf = B + - G ( 1

Experimental procedures

Alloys were obtained from commercial sources in the form

of % in ( 1 .9 cm ) diameter rods. Cylindrical specimens 78 in (1.6 cm) in diameter and % in (1.6 cm) long were prepared

by grinding. Opposite faces were tlat and parallel within 100 X 10"^ in (2.5 nm). Specimens were annealed at a pres- sure of 5x10"*' torr, or less, and cooled in the furnace. Annealing times and temperatures are given in Table 1 , to- gether with chemical compositions (obtained from mill analyses), hardness numbers, microstructures, grain sizes, and mass densities. Hardness, microstructure, and grain size were determined by standard metallurgical methods. Mass

density was determined by Archimedes' method using dis- tilled water as a standard. The specimen holder is shown in Fig.l. The holder was

placed in the ullage of a helium dewar and lowered stepwise

to achieve cooling. Temperatures were monitored with a chromel- constanlan thermocouple contacting the specimen.

Quartz transducers ( 1 0 MHz) were bonded to the specimens

with phenyl salicylate for room-tem'perature measurements

and with a stopcock grease for lower temperatures. In a few

cases, failure of these bonds at very low temperatures re- quired using a silicone lluid (viscosity = 200 000 cP at 25°C)

for bonding.

The authors are with the Cryogenics Division, Institute for Basic Standards, National Bureau of Standards, Boulder, Colorado 80302, USA. ERN and WFW are NRC-NBS Postdoctoral Research

Associates for 1973-74. Received 28 December 1973.

and the transverse modulus is given by

C\ = pv\ = G (2)

Bakelite

cylinder

Thermocouple

Bakelite

cylinder

To co-axial leads

2-6 inductorl Rigidly mounted

4-1 1 inductorj tuning circuit

Brass nut ^ - Brass plunger

Brass txjse

Fig.l Specimen holder

Beryllium-copper spring

Quartz transducer Specimen

Thin-wolled stainless tubing

(electrical ground)

Heat vane (optional)

246 274 CRYOGENICS . MAY 1974
Table 1 Compositions and properties of alloys

Alloy Chemical composition, mill analyses, wt %

Ti Al V So Fe C N

Hardness (DPH No

Mass density Grain size. at 294 K, 1000 g load) average dia, mm g cm '^ Conditions Microstructure

Ti-6A1-4V Balance 6.2 4.0

TI-5A1-2.5Sn Balance 5.5

0.1 0.01 0.01 0.01 320

2.5 0.2 0.07 - 0.02 330

0.005

0.02

4.42

4.47

Annealed

(800° C, 2 h)

Annealed

(816°C, V2 h)

Equiaxed

alpha grains with inter- granular beta Equiaxed

alpha grains

in a mottled

dark-etching

(Kroll's re- agent) matrix

phase, pro- bably acicular alpha

' E

_Z

Table 2 Parameters in equation 3

100 ISO 200

Temperature, K

250 300

Fig. 2 Longitudinal modulus C| = pv\^ of two titanium alloys

Tl-5 AI-2 5 Sn

100 ISO 200

Temperature^ K 2SO 300

Fig. 3 Transverse or shear modulus Ct - C - iivx^ of two titanium

alloys

s. t, Alloy Mode 10^1 N m-2 1011 N m 2 K

Ti-6A1-4V pi/|2 1.688 0.085 213.7

pvt^ 0.467 0.039 183.6

Ti-5A1-2.5Sn pi/|2 1.716 0.085 208.8

pvi^ 0.470 0.051 210.4

Table 3 Temperature derivatives of elastic constants at room temperature (10"^ K"^ )

1 dB 1 dE 1 dG 1 dt^ Alloy

B dT Ed?" G dr w dT

Ti-6A1-4V -1.01 -4.44 -4.90 1.89

Ti-5A1-2.5Sn -0.75 -4.99 -5.55 2.26

Pure Ti* -0.93 -6.63 -7.15 2.80

'Calculated from data in reference 5 using a Voigt— Reuss— Hill average.

Here I'l and t't are the longitudinal and transverse sound- wave velocities, p is the mass density, B is the bulk modulus,

and G is the shear modulus.

Temperature dependences of both C\ and Q were fitted to a semi-theoretical relationship suggested by Varshni

C=(f

e'/^ - 1.

where C , s, and t are adjustable parameters and T is temperature. The value oi C at 7 = 0 K is C°, and s/l is the high temperature limit of the temperature derivative dC'/dr. By invoking an Hinstein oscillator model ot st)lids,

it can be shown (in the absence of electronic effects) that

/ is the Hinstein characteristic temperature. Parameters C°,

.s, and / are given in Table 2. Room temperature values of

the temperature coefficients of the elastic moduli are given

in Table 3.

CRYOGENICS . MAY 1974 275 247
Curves in Figs 2 and 3 are plots of (3) determined by an unweighted least-squares fit of the data. Average percentage differences between measured and curve values are 0.03% and 0.07% for the longitudinal and transverse moduli,

respectively.

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 5, Young's modulus E, the shear modulus G, and Poisson's ratio v. The relationships

among these are

Table 4 Elastic Debye temperatures at T = 0 K

11 1 E 3G 9B

and

E

v = 1 2G

(4)

(5)

These elastic constants were calculated from the moduli

shown in Figs 2 and 3 by the relationships

3Ct I Q - - Q

Q-Ct

B = C\--Ci

3 and

/ Q-2Ct \

(6)

(7)

(8) Q-Cj

The constants E, B, and v are shown in Figs 4-6.

The elastic Debye temperature d is of fundamental import- ance in the lattice properties of solids and is related to the

elastic wave velocities by

where

h (3Np\^l^

4nA

(9)

(10)

Here h is Planck's constant, k is Boltzmann's constant. A' is Avogadro's constant, p is the mass density, and A is the atomic weight. The average velocity is given by

<)>> = 3 -I- 2Ft'^ 1/3(11)

The Debye temperatures for the two alloys at T = 0 K, and

also for pure titanium, are given in Table 4.

Discussion

The elastic properties of both Ti-6A1-4V and Ti-5Al-2.5Sn

behave regularly with respect to temperature. All of the

elastic moduli (C], Q = G, B, E) decrease with increasing

Alloy 6, K

Ti-6A1-4V 418.1

Ti-5A1-2.5Sn 416.9

Pure Ti 425.7*

*Reference 5

temperature, show a relative flatness at low temperatures,

achieve zero slope at 7= 0, and approach linear behaviour

at high temperatures. Poisson's ratio also behaves regularly,

having a positive temperature coefficient. Assuming that the specimens studied are representative of

the two alloys, then conclusions concerning their relative

elastic behaviour can be drawn. Not surprisingly, as shown

in Figs 2—6 and Tables 2 and 3, for most practical purposes

the two alloys are elastically identical. Thus, whether the second alloying element is vanadium or tin, the elastic properties are essentially the same. As is well known, titanium undergoes a cph to bcc (a to (3) crystal-structure transition when heated to 1 155 K. Fisher and Renken ^ found a large temperature dependence of the

£-66 shear modulus in titanium

1 dC66

11.93 X lO--* K->

A small value of Cg^ is expected from the shear mechanism

of a cph to bcc transition proposed by Burgers.^ For hexagonal symmetry, the shear modulus G is an approximate

average of the two single-crystal shear constants c^^ and C44. Thus, from the smaller values of (I /G){dG/dT) in Table 3 one can infer that both alloys have higher cph to bcc trans- ition temperatures than pure titanium. That is, the a-titanium phase is stabilized in the alloys. As shown in Hansen and Anderko,'' aluminium is a strong stabilizer of a

titanium, while both vanadium and tin are moderate

stabilizers of j3 titanium.

1.264

Fig.4

100 ISO 200

Temperature^ K

Young's modulus of two titanium alloys

250 300

248 276 CRYOGENICS . MAY 1974
Table 5 Comparison of present and previously reported results; room temperature values; units of 10^ ^ N m'^ except

I' (dimensionless)

Source

Ti-5A1-2.5Sn

G B

Ti-6A1-4V

G B

Reference 12 Reference 8 Reference 9 Reference 10 Reference 1

1

Present

1.14 0.96-1.17

1.07 1.10

1.11 0.420

1.07

1.07 0.327

1.11

1.07- 1.10

1.03

1.1

1

1.21

0.421

0.421

0.420

1.06

1.13

1.05

0.325

0.323

I.098 1092

l,086-

'^E I080- z

=0 I074|-

5 I 068 ^

1 I062-

E^ 1056-

I 050-

1044-

038

TI-5AI-2 5 Sn

TI-6 AI-4V

o 50 lOO ISO 200 2 SO 300

Temperature, K

Fig. 5 Bulk modulus (reciprocal compressibility) of two titanium

allovs O 338r 0334

0 298

Fig.6

lOO ISO 200

Temperature, K

Poisson's ratio of two titanium alloys

300

Despite the extensive use of these alloys, very little elastic data exist for them. Most information has appeared in engineering reports and is summarized in references 8-12.

For comparison, tiie room temperature values of E, G, B. and V are given in Table 5. Generally good agreement is observed between previous and present results. The clastic moduli of both alloys are several percent higher

than those reported for commercial-grade titanium.'^ But

they are lower by several percent than the quasi-isotropic

moduli obtained from a Voigt— Reuss-Hill average of the data obtained from zone-refmed single crystals.^ The elastic Debye temperatures of the alloys are also about 2% lower

than that calculated from the single-crystal elastic data of

titanium. The relatively low elastic anisotropy of titanium does not allow for a large error in the Voigt-Reuss— Hill averages. Thus, existing data suggest a strong impurity

effect on the elastic constants of titanium, perhaps due to

interstitial impurities.

Finally, 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. Conversion formulas are given in Landau

and Lifshitz,'^ for example. For titanium at room

temperature

0.001 5s - Bj

= 0.009

Gc — G-r 0.005 and -^0

where subscripts S and T denote adiabatic and isothermal,

respectively.

This work was supported in part by the Advanced Research

Projects Agency of the Department of Defense and was monitored by the Cryogenics Division, NBS, Boulder under

contract No 2569.

References

1 McSkimin, H. J. J Acoiist Soc Amcr 3i (\96\) 12

2 McSkimin, H. J., Andreatch, P. J Acoiist Soc Attwr 34 (1962) 609

3 VarshnI, Y. P. P/n s B2 (1970) 3952

4 Debye, P. Ann Phys (Lcipzj 39 ( 1 9 1 2) 789

5 Fisher, E. S., Renken, C. J. Phys Rev 1 35 ( 1 964) A482

6 Burgers, W. G. Physica 1 (1934) 561

7 Hansen, M., Anderko. K. Constitulicm ot Binary Alloys (McGraw-Hill, New York. 1958)

8 Metals Handbook, 8th eiln, Vol 1 : Properties and Selection of

Metals (Anicr Soc Metals, Metals Park, Ohio, 1961

)

9 Cryogenic Materials Data Handbook, AFML Rep No ML TDR-64-280 (PB 171809, revised, 1964)

10 NFRVA Proiirani Materials Properties Data Book, Vol 1

:

Introduction and light Metal Alloys (Aerojet Nuclear Systems, Sacramento, Calif, 1970)

I 1 Hanlein, S. L., Hinckley, W. M., Stecher, F. P. NOLTR Report 70-141 (1970)

12 Fahey,N.H. W AL Tech Rep No TR 1 1 8. 1 / 1 ( I 960)

1 3 Landau, L. D., Lifshitz, E. M. Theory of tiasticity (Pcruamon, London 1959) 17

CRYOGENICS . MAY 1974 277

249
Nba-l 14A (REV. 7-73)

U.S. DEPT. OF COMM. BIBLIOGRAPHIC DATA

SHEET

1. PUBLICATION OR REPORT NO. NBSIR74- 393

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

by Reed, Durcholz, Fickett, Giarratano, Hust, Kasen

5. Publication Date October 1974

6. Performing Organization Code

275. 03

7. AUTHOR(S) Ledbetter, Mikesell, Naimon, Schramm, Sparks

Tobler^and Weston

8. Performing Organ. Report No.

9. PERFORMING ORGANIZATION NAME AND ADDRESS

NATIONAL BUREAU OF STANDARDS , Boulder Labs.

DEPARTMENT OF COMMERCE

WASHINGTON, D.C. 20234

10. Project/Task/Work Unit No. 2756530

11. Contract/Grant No. ARPA Order No. 2569

12, Sponsoring Organization Name and Complete Address (Street, City, State, ZIP)

Advanced Research Projects Agency

Department of Defense

Washington, D. C.

13. Type of Report & Period Covered

14. Sponsoring Agency Code

ARPA

15. SUPPLEMENTARY NOTES

16. ABSTRACT (A 200-word or less factual summary of most sigrjificant information. If document includes a significant

bibliography or literature survey, mention it here.) Results of six months of study on Materials Research in Support of Superconducting Machinery (April through September, 1974) are reported to the sponsor,

the Advanced Research Projects Agency of the U.S. Department of Defense. The

report is divided into five sections: thermal conductivity, magnetothermal

conductivity, fatigue and fracture-toughness properties, properties of advanced

composites, and elastic properties. The temperature range 4 to 300 K is covered

by the study. Materials studied are either being used or are candidates for use in

superconducting machinery and include: aluminum alloys , composites, inconels,

OFHC copper, stainless steels, and titanium alloys.

Special results of the study include: fracture -toughness and fatigue-crack growthrate data for AISI 310, Ti-5 Al-2.5 Sn, A286 stainless steel, and Inconel 750 at 4,

76, and 300 K; initial reports of tensile testing of composites at 4 K; a second review

paper on advanced- compos ite low-temperature behavior; magnetothermal conductivity

measurements on OFHC copper and Inconel 718 indicate that, in a magnetic field, the

thermal conductivity may change by 100 percent; anomalous low-temperature elastic

behavior of AISI 300 series stainles s - steel alloys; and thermal conductivity data for

OFHC copper, Inconel 750, and Inco Low-Expansion-Alloy. These data provide consi

derable insight into material characteristics at extremely low temperatures, assisting

17. KEY WORDS i in material selection and efficient design.

name; separated by semicolons

)

elastic properties*

Composites;^fracture; liquid helium; mechanical properties; structural materials;

superconducting machinery; thermal conductivity.

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|~y For Official Distribution. Do Not Release to NTIS

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