NBSIR 76-839

SEMI-ANNUAL REPORT ON MATERIALS RESEARCH

IN SUPPORT OF SUPERCONDUCTING MACHINERY

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

D. T. Read, R. E. Schramm, L. L. Sparks, and R. L. Tobler

Cryogenics Division

Institute for Basic Standards

National Bureau of Standards

Boulder, Colorado 80302

April 1976

Prepared for : Advanced Research Projects Agency

1400 Wilson Boulevard

Arlington, Virginia 22209
,1

I
NBSIR 76-839

SEMI-ANNUAL REPORT ON MATERIALS RESEARCH

IN SUPPORT OF SUPERCONDUCTING MACHINERY

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

D. T. Read, R.E.Schramm, L. L. Sparks, and R. L. Tobler

Cryogenics Division

Institute for Basic Standards

National Bureau of Standards

Boulder, Colorado 80302

April 1975

Prepared for : Advanced Research Projects Agency

1400 Wilson Boulevard

Arlington, Virginia 22209

U.S. DEPARTMENT OF COMMERCE, Eliiot L Richardson, Secretary

James A. Baker. III. Under Secretary

Dr. Betsy AncKer-Johnson . Assistant Secretary for Science and Technology

NATIONAL BUREAU OF STANDARDS, Ernest Ambler, Acting Director

SEMI-AXNTAL REPORT ON MATERIALS RESEARCH

IN SUPPORT OF SUPERCONDUCTING MACHINERY

Sponsored by

Advanced Research Projects Agency

ARPA Order No. 2569

Program Code 4D10

August 10, 1975 - October 1, 1976

Program Director

Dr. E. C. van Reuth

Materials Sciences

Advanced Research Projects Agency

1400 Wilson Boulevard

Arlington, Virginia 22209

Program ^^anager

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.

lii
Abstract

Results are reported of a six-month study, ending March 1976, on candidate materials

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

properties, fatigue resistance and fracture toughness, magnetothermal conductivity, and

thermal conductivity. Material properties were studied over the temperature range 4 to 300 K. Materials studied include: aluminum alloys 1100, 2014, 2219; a nickel-chromium-iron

alloy; iron-47.5 nickel; and the composite materials boron/aluminum, boron/epoxy, S-glass/

epoxy; graphite/epoxy . Some notable results of the study are: first reports of compressive

mechanical testing on composite materials at 4 K; regular temperature behavior of the elastic constants of aluminum 2014 and 2219 and of iron-47.5 nickel, which is magnetic; none of

the mechanical properties of the nickel-chromium-iron alloy tested were affected deleteriously by cryogenic temperatures; in aluminum alloy 2219, Jj^ ^^'^ ^Jc equivalent

because of sub-critical crack extension; both electrical and thermal conductivities of

aluminum alloy 1100 are reduced by magnetic fields.

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

of Defense.

Keywords: Aluminum alloys; composites; elastic properties; engineering materials; fatigue;

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

superconducting machinery; thermal conductivity.

iv
Report Contents

Page

ADVANCED COMPOSITES

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

Reprint: Mechanical and Thermal Properties of Filamentary-Reinforced

Structural Composites at Cryogenic Temperatures. 2: Advanced

Composites; M. B. Kasen; Cryogenics 15 (1915) 701-22 19

ELASTIC PROPERTIES OF ENGINEERING MATERIALS AT CRYOGENIC TEMPERATURES

H, M. Ledbetter and D. T. Read 41

Low-Temperature Elastic Constants of Precipitation-Hardened Aluminum

Alloys 2014 and 2219

D. T. Read and H. M. Ledbetter 43

Low-Temperature Elastic Properties of an Iron-47.5 Nickel Alloy

H. M. Ledbetter and D. T. Read 58

Reprint: Low-Temperature Elastic Properties of a Nickel-ChromiumIron Alloy; W. F. Weston and H. M. Ledbetter; Mater. Sci. Eng 2Q_ (1975) 287-90 66

Reprint: Low-Temperature Elastic Properties of Four Wrought and

Annealed Aluminum Alloys; E. R. Naimon, H. M. Ledbetter and

W. F. Weston; J. Mater. Sci. 10^ (1975) 1309-16 70

Reprint: Low-Temperature Elastic Constants of a Superconducting

Coil Composite; W. F. Weston; J. Appl. Phys . 46^ (1975) 4458-65 78

FRACTURE MECHANICS PARAMETERS OF ENGINEERING MATERIALS AT CRYOGENIC TEMPERATURES

R. L. Tobler, D. T. Read, and R. P. Reed 87

Low-Temperature Effects on the Fracture Behavior of Inconel 718

R. L. Tobler 89'

Effects of Specimen Thickness on Fracture Toughness of Some

Aluminum Alloys

D. T. Read and R. P. Reed . , . . . 110

MAGNETOTHERMAL CONDUCTIVITY

L. L. Sparks 139

THERMAL CONDUCTIVITY

J. G. Hust 149

V
Disclaimer

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

materials for discussion and examination with regard to application in superconducting

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

or equipment is necessarily the best available for the purpose.

vi
SEMI-ANNUAL REPORT ON MATERIALS RESEARCH

IN SUPPORT OF SUPERCONDUCTING MACHINERY

ADVANCED COMPOSITES

R. E. Schramm and M. B. Kasen

Cryogenics Division

Institute for Basic Standards

National Bureau of Standards

Boulder, Colorado 80302

April 1976
SuiTiinary; Advanced Composites

Phase II of the experimental program has been completed. This report

contains 2 95 K, 76 K and 4 K uniaxial compressive laminate mechanical property data on the following commercial composites: 5.6 mil boron/6061 aluminum, 5.6 mil boron/5505 epoxy*, S-901 glass/NASA Resin 2 epoxy, and type

AS graphite/NASA Resin 2 epoxy. A reprint of a review of the literature on

the mechanical and thermal properties of advanced composites at cryogenic

temperatures is also included.

Contents; Advanced Composites

Page

1.0 Review 3

2.0 Phase II (Continued): Static Mechanical Properties

of Uniaxial Composites at Cryogenic Temperatures 3

2.1 Introduction 3

2.2 Experimental Procedures 4

2.3 Results and Discussion 5

2.3.1 Compressive Modulus 5

2.3.2 Compressive Strength 5

2.4 Conclusions 6

3 . 0 Future Work 7

4 . 0 References ^

5.0 List of Figures 1

6.0 List of Tables 8 Appendix I: Calculation of Compressive Moduli 18

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

properties of filamentary-reinforced

structural composites at cryogenic

temperatures. II: Advanced

composites," Cryogenics 15 (12),

pp. 701-722 (1975) 19

* The use of trade names or designations in this paper is essential to

proper understanding of the work presented. Their use in no way implies

approval, endorsement, or recommendation by NBS

.

2
1.0 Review

Work on this project was initiated by preparing state-of-the-art reviews

on the mechanical and thermal properties of structural composites at cryogenic

temperatures. A general review has been published ( 1 ) . More comprehensive

reviews of glass-reinforced composites and of advanced-fiber (high-modulus)

composites have been published (2 , 3) . A reprint of Ref. 3 is included in

the present report. Reprints of Refs. 1 and 2 have appeared in a previous

Semi-Annual Report (4).

The reviews suggested that boron-aluminum, boron-epoxy, graphite-epoxy

,

glass-epoxy and Kevlar 4 9-epoxy composites warranted further study to characterize key static mechanical properties at cryogenic temperatures. Key

properties are those required for a prediction of strength or stiffness

limits in complex crossply layups using macromechanical composite theory.

The properties are obtained from uniaxial composite laminates , and consist of

strengths and elastic moduli in tension and compression in the longitudinal

and transverse directions plus in-plane elastic shear moduli.

Phase I of the experimental work was the development of apparatus and

procedures for obtaining these data at cryogenic temperatures. Results have

been reported previously ( 5 , 6 ) . Phase II encompassed the static characterization work. The tensile and in-plane shear portion of this Phase was reported

previously ( 4 ) . The static compressive mechanical properties of these composites are presented and discussed in the present report.

In Phase III of this work, the performance of boron-epoxy and boronaluminum composites in tensile-tensile fatigue will be examined at cryogenic

temperatures. This work is presently under way and will be reported subsequently .

2.0 Phase II (Continued) : Static Mechanical Properties

of Uniaxial Composites at Cryogenic Temperatur"es

2.1 Introduction

This report presents compressive static mechanical property data for the

composites listed in Table 1, all of which had previously been characterized

in tension and in-plane shear. The boron-reinforced composites were fabricated from state-of-the-art commercial preimpregnated tape materials, not

optimized for cryogenic use. The cryogenic properties reported for these

materials in the present work are believed representative of the properties

of boron-reinforced aluminum or epoxy composites as a class, independent of

the specific manufacturer.

NASA Resin 2 is a non-proprietary formulation, developed for cryogenic

filament-wound pressure vessles(7). It consists of Epon 828/DSA/Empol 1040/

BDMA in proportions of 100/115.9/20/1 by weight.

The Kevlar 49/NASA Resin 2 composite, previously included in the tensile evaluation, was not included in the compressive test program. The very

low transverse strength of the material obtained for this program suggested

that the composite was not representative of good production quality.

Key compressive static mechanical properties include uniaxial longitudinal and transverse ultimate strengths, elastic moduli, and ultimate failure

strains. Where possible, the stresses at the proportional limit and the

0.2% offset yield strengths were also determined.

The present data have been obtained from rod or bar type specimens

3
designed to fail in approximately 4 5° shear. This is believed to best

approximate compressive failure in bulk composite structures in the absence

of significant column bending or end brooming. The reader should be aware

that compressive properties obtained by the sandwich beam method may differ

somewhat from those reported in the present work.

2.2 Experimental Procedures

The compression fixture shown in Figure 1 converts pull rod . separation

into compression by means of interlocking yokes. Specimen (A) with its end

caps is inserted into compression blocks (B) . Alignment is maintained by

sleeve (C) , which slips around the blocks. The compression blocks are anodized to minimize friction. Additionally, M0S2 lubricant sprayed on the

fixture during assembly serves as both a release agent for the end caps and

as a lubricant between the compression blocks and alignment tube. Extraneous

frictional forces are typically 5-10 N (1-2 lb)

.

The split compression blocks are joined by stainless steel bolts,

facilitating removal of the end caps after testing. Fixture dimensions and

construction materials are detailed in Figure 2.

The compression fixture was designed to interchange with the tensile

fixture in the cryostat previously described ( 5) . Helium consumption was typically 5-6 liters per specimen test.

The desired 45° shear failure mode was reliably obtained after several

iterations of specimen design. The two final configurations are shown in

Figure 3. Specimens of square cross section, embedded in 6061-T6 aluminum

end capSf proved satisfactory for most materials and orientations (Figure 3 (a)

^

However, the very high compressive strengths developed by the boron-reinforced

materials when tested along the fiber direction required round specimens and

stainless steel end caps (Figure 3 (b) ) . Both specimen types were diamond

ground to finished dimensions of Figure 3.

A longitudinally-slit tube, having an i.d. matching the o.d. of the end

caps, was used to align and support the specimens during room temperature

curing of the epoxy used to bond the specimens into the end caps. Several

conventional types of epoxy proved satisfactory for this application.

| J All polymer-matrix composites were environmentally conditioned for a [ minimum of 40 hours at 23 + 1° C in 50 + 10% relative humidity prior to

testing. The boron/6061 aluminum was tested in the as-fabricated condition

(F temper) without additional conditioning.

A universal testing machine was used for all tests. The crosshead speed was 0.10-0.13 cm min~^ (0.04-0.05 in min~l) , providing a strain rate of 0.06-

0.08 min-1.

Specimens of square cross-section were instrumented with a single longitudinal strain gage. On transverse specimens, the gage was always on the

side parallel to the fiber reinforcement. The initial tangent modulus was obtained from load-displacement traces recorded at high sensitivity on an x-y plotter; maximum stress was limited to less than 25% of the ultimate

strength. Specimens were subsequently compressed to failure while recording

load and strain at lower sensitivity. The strain gage typically failed

prior to specimen fracture. However, simultaneous load-time traces on a strip-chart recorder permitted calculation of overall failure strain.

Strain gages were not used with specimens of round cross section.

However, the compressive moduli of these specimens were calculated from the

load-time trace of the strip chart recorder using the procedure described in

Appendix I .

A
2.3 Results and Discussion

Individual specimen data appear in Tables 2-5. Table 6 presents average

values, omitting data from invalid fractures when calculating ultimate strain

and strength. The inaccuracy of the data is estimated at + 3% for ultimate

strength, + 10% for modulus and ultimate strain, and + 20%~for proportional

limits and yield strengths.

The maxima of the compressive ultimate strengths and elongations are presented in Table 7. It is probable that the values in this Table are a better reflection of the true ultimate properties of the bulk composites than

are the average values of Table 6. Particularly when testing in the uniaxial

longitudinal direction, values obtained from relatively small cross section

rod specimens are probably lower due to premature failure in other than pure

compression.

The desired shear mode of compressive failure is evident in the fractures of the various materials illustrated in Figure 4.

2.3.1 Compressive Modulus

The compressive moduli of the boron-epoxy, glass-epoxy and graphiteepoxy materials are very close to their tensile moduli in both the longitudinal and transverse directions. In contrast, the longitudinal compressive

moduli of the boron-aluminum composite was computed to be 30-80% higher than

the tensile value, while the transverse compressive moduli was computed to

be 10-30% less than the corresponding tensile value.

The authors do not believe that the compressive moduli calculated for

the boron-aluminum composite represent valid material properties. An average

longitudinal compressive modulus of 359 GN/m^ (52.1 X 10^ psi) is

hardly realistic for a composite reinforced with 50 v/o of boron

which itself has a modulus of about 379 GN/m^ (55 X 10^ psi) . Furthermore,

preliminary studies of the elastic constants of the same boron-aluminum

composite using dynamic resonance techniques indicate that the average of

the tensile and compressive moduli is about 234 GN/m^ (34 X 10^ psi) , with

a small temperature sensitivity on cooling to cryogenic temperatures ( 8 ) . It

therefore appears that the method used to obtain compressive moduli in the

present work is invalid for boron-aluminum composites. We do not know the

reason for this discrepancy. However, until the problem is resolved, the

authors suggest that tensile moduli values be used to approximate compressive moduli when working with boron-aluminum composites.

The longitudinal compressive moduli of the boron-epoxy and glass-epoxy

composites increased about 20% on cooling to 4 K, while that of the graphiteepoxy composite remained relatively unchanged. A much larger temperature

dependence of modulus was observed in the transverse direction, with increases

of 67%, 108% and 178% being observed for the graphite-epoxy , boron-epoxy,

and glass-epoxy, respectively. Similar temperature dependences had been

previously observed for the tensile moduli of these materials.

2.3.2 Compressive Strength

Present data suggest that boron-reinforced composites are capable of

providing longitudinal compressive strengths approaching 3.4 GN/m^ (50 X

10"* psi) at 4 K in the absence of column bending. This is twice the tensile

strength at that temperature. Conversely, the 4 K longitudinal compressive

strengths of the glass and graphite-epoxy composites were 30-50% lower than

their 4 K longitudinal tensile strengths. The superior ability of the boron

fibers to sustain compressive loads is clear.

The transverse compressive strength of all four composite types are substantially higher than their transverse tensile strengths. This is

5
particularly noticeable in the graphite-epoxy composite.

The compressive strength of the boron-reinforced composites in the longitudinal direction appears to be relatively independent of temperature,

although present data are incomplete. However, the longitudinal compressive

strengths of the glass-epoxy material is very temperature dependent, increasing by 180% on cooling. A 25% increase in strength was observed for the

graphite-epoxy on cooling. In the transverse direction, the compressive

ultimate strength of the boron-aluminum, boron-epoxy, and glass-epoxy increased

by 100-170% on cooling. However, cooling to cryogenic temperatures decreased

the transverse compressive strength of the graphite-epoxy composite by about

30%.

2.4 Conclusions

Commercial, state-of-the-art boronr6061 aluminum and boron-epoxy composites have excellent static compressive properties at cryogenic temperatures.

The compressive strength is particularly outstanding, approaching 3.4 GN/m^

(50 X 10^ psi) in the uniaxial longitudinal direction at 4 K. Transverse

compressive strengths of 620 MN/m^ (90 X 10-^ psi) for boron-aluminum and 427

MN/m2 (62 X 10"^ psi) for boron-epoxy at 4 K are substantially higher than in

other composite systems. The compressive strength properties of the boronreinforced materials are superior to the already high tensile strength properties at all cryogenic temperatures. With a modulus of 206-240 GN/m^

(30-35 X 10^ psi), the boron-reinforced composites appear to be excellent

cryogenic structural materials.

At 4 K, NASA Resin 2 epoxy reinforced with S-901 glass fibers has about

half the compression strength of the boron materials in both the longitudinal

and transverse directions. The longitudinal compressive strength of this

composite is 30-60% lower than the corresponding tensile strength, while the

transverse compression strength is 100-200% greater than in tension. Compressive moduli increase with cooling, attaining about 62 GN/m^ (9 x 10^ psi)

longitudinally and about 31 GN/m^ (4.5 X 10^ psi) in the transverse direction

at 4 K.

NASA Resin 2 reinforced with Type AS graphite fiber had the lowest com- pressive strength of the composites tested. Uniaxial longitudinal strengths

at 4 K were about 25% of that of the boron materials, while transverse

strengths averaged about 130 MN/m^ (19 X 10-^ psi) . The longitudinal compression modulus of about 117 GN/m^ (17 X 10^ psi) was about the same as the

tensile modulus.

The comparatively poor performance of the graphite-epoxy composite in

the present work should not be construed as indicating a general inferiority

of this type of composite for cryogenic applications. The data suggest that

an incompatibility may exist between the fiber and matrix in the AS/NASA

Resin 2 composite selected for this study.

Present data indicate that the compressive and tensile moduli of the

boron-epoxy, graphite-epoxy and glass-epoxy composites are substantially the same from 295 K to 4 K. In view of the experimental difficulty, there

appears to be little justification for a separate determination of compressive moduli if tensile moduli are available.

We conclude from this study that the compressive moduli of uniaxial

boron-6061 aluminum composites cannot be reliably measured by the procedures

used in this study, even though such procedures were satisfactory for the

epoxy-matrix composites.

6
3 . 0 Future Work

Phase III of this program will examine the extent to which cryogenic

temperatures affect the wear-out rate of boron-aluminum and boron-epoxy com- posites under tension-tension fatigue. Criteria will be changes in specimen

modulus and damping. This study is presently under way using composites of

[0/+ 45/0] s orientation and stacking.

4.0 References

1. Kasen, M. B., "Properties of Filamentary-Reinforced Composites at Cryogenic Temperatures," Composite Reliability , ASTM STP 580, American

Society for Testing and Materials, 1975, pp. 586-611.

2. Kasen, M. B. , "Mechanical and Thermal Properties of Filamentary-Reinforced

Structural Composites at Cryogenic Temperatures. I: Glass-Reinforced

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

3. Kasen, M. B. , "Mechanical and Thermal Properties of Filamentary-Reinforced

Structural Composites at Cryogenic Temperatures. II: Advanced Composites,"

Cryogenics, 15 (12), 1975, pp. 701-722.

4. Schramm, R. E. and Kasen, M. B., "Advanced Composites," in Semi-Annual

Technical Reports on Materials Research in Support of Superconducting

Machinery - IV ; R. P. Reed, A. F. Clark, E. C. van Reuth (Eds.); Nat.

Bur. Stds., Boulder, CO; October 1975; ADA019230.

5. Kasen, M. B., "Advanced Composites," in Semi-Annual Technical Reports

on Materials Research in Support of Superconducting Machinery ; R. P.

Reed, A. F. Clark, E. C. van Reuth (Eds.); Nat. Bur. Stds., Boulder, CO;

March 1974; AD780596.

6. Kasen, M. B. and Schramm, R. E., "Advanced Composites," in Semi-Annual

Technical Reports on Materials Research in Support of Superconducting

Machinery - II ; R. P. Reed, A. F. Clark, E. C. van Reuth (Eds.); Nat.

Bur. Stds., Boulder, CO; October 1974; ADA004586.

7. Soffer, L. M. and Molho, R. , "Cryogenic Resins for Glass Filament-wound

Composites," NASA CR-72114 (Final), 1967, N67-25076.

8. Read, D. T. and Ledbetter, H. M. , NBS-Boulder, personal communication

(March 1976)

,

5.0 List of Figures

1. Fixture for compression testing of composites at cryogenic temperatures.

(a) Specimen with end caps

(b) Anodized aluminum compression blocks

(c) Aluminum alignment sleeve

2. Drawing of fixture for compression testing of composites at cryogenic

temperatures

.

3. Specimen configurations used in determining static compressive properties

of composites at cryogenic temperatures.

(a) Square specimen for glass-epoxy, graphite-epoxy and transverse

boron-aluminum and boron-epoxy

(b) Round specimen for longitudinal boron-aluminum and boron-epoxy

7
4. Typical composite compressive fractures.

(a) Glass-epoxy, longitudinal, 295 K

(b) Glass-epoxy, transverse, 295 K

(c) Glass-epoxy, longitudinal, 4 K

(d) Boron-aluminum, longitudinal, 4 K

(e) Boron-epoxy, longitudinal, 4 K

(f) Boron-epoxy, transverse, 4 K

(g) Graphite-epoxy , longitudinal, 295 K

(h) Graphite-epoxy, longitudinal, 4 K

6.0 List of Tables

1. Composites Tested in Compression

2. Compression Properties of 5.6-mil Bbron-6061 Aluminum Composite

(Individual Specimens)

3. Compression Properties of 5.6-mil Boron-5505 Epoxy Composite

(Individual Specimens)

4. Compression Properties of S Glass-NASA Resin 2 Composite

(Individual Specimens)

5. Compression Properties of Type AS Graphite-NASA Resin 2 Composite

(Individual Specimens)

6. Compressive Properties of Composites (Averages of Specimens Tested)

7. Compressive Properties of Composites (Maxima of Ultimate Strength Values)

8
Fixture for compression testing of composites at cryogenic temperatures.

(a) Specimen with end caps

(b) Anodized aluminum compression blocks

(c) Aluminum alignment sleeve

8

16

100(L)

3.56x5.0x42mm aluminum

strip. Epoxy in place, both

sides of specimen.

Machine round.

12.7

^Mill slot 3.57x42mm long

-3.56°

50(T)

n

3.18D-

#304 Stainless Steel

3.19 mm drilled hole,

42 mm deep. Epoxy

specimen as shown.

6061-T6

Aluminum

lJ_

"f 100

16

_1 118

a

All dimensions shown in mm.

Fig Specimen configurations used in determining static compressive properties of composites at cryogenic temperatures

(a) Square specimens for glass-epoxy, graphite-epoxy

and transverse boron-aluminum and boron-epoxy.

(b) Round specimen for longitudinal boron-aluminum

and boron-epoxy.

11
1 fcrAX.Wtu.

1.

-

AT- >-l.\t.-

I

I sin u5i-

—

VAi

1. ®

(5.)

IS.«.M.V\»>V-W

12
Table 1. Composites Tested in Comoression

Average

Nominal Fiber

Plate Thickness Ply Thickness Volume Composite Type (cm) (in) (cm) (in) Percent*

5.6 mil Boron-6061 Aluminum

(condition F)

0 373 0 147 0 0175 0 0069 47

5.6 mil Boron-5505 Epoxy 0 345 0 137 0 0170 0 0067 52 S-901 Glass-NASA Resin 2 Epoxy 0 462 0 182 0 0206 0 0081 66

Type AS Graphite-NASA Resin 2 Epoxy 0 366 0 144 0 0193 0 0076 64

* Plate-to-plate variation was 2%.

Table 2. Compressive Properties of 5.6 mil Boron-6061 Aluminum Comooslte^ (Individual Specimens)

Temperature Young's Modulus •

I Prooortlonal Limit 0.2X Yield Strength, a'^^ Ultimate Strength, 0=" Ultimate

(K) (10* f/m2) (10" psi) (108 N/m2) (10^ psi) (108 N/„2) (103 psi) (108 N/m2) (10^ psi) Strain, £<^" (%)

Longitudinal (0°)

295 221 32 0 19 2 278 1 8 295 293 42 5 15^

76 311 45 1 21 7 m. 76b 407 '

* 59 0

= * 28 7 416 0

5'-

4b 335 48 6" 28 5 413 0

6^= 4b 337 48 8

= 26 0 377 0

5'= 4b 302*^ 43 8"= 32 6 473 0 7"

Trai averse (90°)

295"* 150 21 2 0.588 8 53 1.43 20.8 2 60 37 7 0 4

= 295 186 27 0 0.545 7 97 1.43 20.7 2 54 36 8 0

4== 295 107 15 5 0.594 8 61 1.03 15.0 2 59 37 5 0 5"

76 101 14 7 0.833 12 1 1.95 28.4 4 56 65 2 0 6"

76 89 * 12 9 * 0.818 11 9 2.06 30.0 4 45 64 6 0 8'

76 101 14 6 0.967 14 0 1.99 28.9 4 50 65 3 0 9

=

4 123 17 8 1.13 16 4 2.82 40.9 6 56 95 1 0

7'= 4 90 13 0 1.11 16 0 2.44 35.3 6 29 91 2 1 2"

4 90 13 1 1.01 14 6 2.39 34^7 5 83 84 5 1 0

= —

^Square specimens except as marked. bRound specimen. 'Estimated from strip-chart recording.

''this is the only boron-aluminum specimen that was load cycled.

Underlined values represent invalid fractures.

* See text discussion on validity of these values.

13
Temperature Elastic Modulus, E Proportional Limit

,

Pl 0.2% Yield Strength, a'^'' Ultimate Strength, Ultimate

(K) (109 N/in2) (10b psi) (10^ N/[„2) (10-* psi) (IQB N/m2) (103 psi) (IO8 N/„2) (10^ psi) Strain,

(%)

Ltudinal (0°)

295^ 186 27 0 11 8 171 0._7 190 27 6 12 9 186 0. 8 295b 260 37 7 14__1 204" £±

763 225 32 6 22 8 331 1 1 763 234 33 9 lO" 2gT 0..8 43 239 . 34 13 Q 188

4'= 257

^

37 25 "9 375 0 233 ^ 33 19 2 278 0 4c 291 42 2 36 6 531 0

6'*

Traiisverse (0°)

295a 20.9 3 03 0. 395 5 7 1.78 25.8 2 27 32 3 1. 6 295a 13.3 2 02 0.975 14 1 1 41 20 5 1 6 295b 20.0 2 90 0.715 10 4 1.79 25.9 2 10 30 5 1. 2 295b .22.4 3 24 0.790 11 4 1.80 26.1 1 96 28 4 1 4

76^ 38.7 5 62 2.31 33 6 4 23 61 4 1 2 76^ 39.0 5 64 1.9B 28 6 3 84 55 6 0 5 763 37.7 5 46 1.66 24 1 1 80 26 1 0 5

76'' 39.5 5 74 2.72 39 5 4 56 66 2 1 5 76b 41.4 5 99 3.38 49 0 4 36 63 1 1. 1

4b 42.8 6 20 4 24 61 4 1 1 4b 37.8 5 49 1 67 24 2 0 4 4b 41.3 5 98 or 65 5 1 1 4b 39.4 5 71 4.08 59 2 0 9

^ Rectangular specimen.

^ Square specimen.

^ Round specimen. Estimated from strip-chart recording.

Underlined values represent invalid fractures.

Table 4. Compressive Properties of S Glass-NASA Resin 2 Composite (Individual Soecimens)

Temperature Elastic Modulus Proportional Limit 0.2% Yield Strength, a'^^ Ultimate Strength, 0^" Ultimate

(K) (109 N/n2) (106 psi) (108 N/m2)

1

(10^ psi) (10° N/m2) (10^ psi) (10» N/„2) (103 psi) Strain,

(%)

Longi tudinal (0°)

295 40 4 5 87 1.36 19 8 4 72 68 5 1 2 295 59 2 8 44 0.93 13 5 5 62 81 6 1 3 295 53 2 7 71 3.14 45 6 4 75 68 8 0 9

76 69 7 10 1 1.78 25 8 7 32 106 16 5 239 3 0 76 64 2 9 3 3.24 47 0 9 87 143 13 4 195 2 5 76 72 5 10 5 3.07 44 4 10 3 150 1 5 76 48 6 7 0 4.47 64 8 7 03 102 14 0 203 2 53

4 53 8 7 81 2.44 35 3 9 08 132 14 1 204 1 5 4 63 0 9 13 3.81 55 2 7 56 110 15 4 224 2 63

4 64 7 9 38 3.11 45 1 13 1 190 1 83

Trans verse (90°)

295 11 7 1 70 0.550 7 97 0.733 10.6 0 962 13 9 3 4 295 12 3 1 78 0.310 5 51 0.753 10.9 0 935 13 6 2 6 295 9 8 1 42 0.422 6 12 0. 703 10.2 1 052 15 3 3 4

76 19 8 2 87 0.92 13 4 2 35 34 0 0 7 76 26 0 3 77 1.31 19.0 2 «4 41 2 1 1 76 22 7 3 29 1.19 17.2 3 11 45 1 0 9

4 27 2 3 95 2 76 40 0 1 1

4 39 4 5 71 2 99 43 4 0 8

4 27 3 3 96 2 13 30 9 0 7

Estimated from strip chart recording

.

All square (pacimens

14
Table 5. Compressive Properties of Type AS Graphite-NASA Resin 2 ComDosite (IndiviHual Specimens)

Temperature Elastic Ml (109 N/„2

}dulus

(10^ psi) ProDortlonal

(10^ N/„2) Limit,

(10^ psi) 0.2% Yield St (I08 N/m2) rength, 0*^^ (103 psl) Ultimate Stn

(108 N/m2)

jngth, a"^" (103 psi) Ultimate

Strain e'^" d)

Longi tudinal (0°)

295 137 19 8 2 25 32 6 5 61 81 2 0.5

295 125 18 1 2 66 38 6 4 96 71 9 0.4

295 125 18 2 3 55 51 5 5 38 78 0 0.4

76 115 16 7 4 54 65 9 7 95 115 0.8

76 97 14 1 4 75 69 0.5

76 110 16 0 3 95 57 3 6 72 97 0.7

76 141 20 4 5 04 73 1 12 29 178 1.0

4 109^ 15 8* 8 83 128 0.7*

4 124 18 0 3 70 53 7 5 87 85 0.6

4 125 18 2 3 77 54 7 5 99 87 0.8

Transverse (90°)

295 9 09* 1 32* 0 90* 13 0* 1 04 15 1 0.8*

295 9 99a 1 45* 0 583 8 4a 0 91 13 2 0.8a

295 9 91* 1 44* 0 71 10 3 0.6*

76 16 6 2 41 1 16 16 8 1 51 21 9 1.1

76 13 1 1 89 0 80 11 5 1 25 18 1 1.0

76 16 1 2 34 0 90 13 1 1 33 19 3 0.9

4 18 8 2 72 0 660 9 58 1 00 14 4 0.5

4 14 4 2 09 0 723 10 48 1 50 21 8 1.0

4 15 3 2 22 0 647 9 39 1 43 20 7 0.9

* Estimated from strip-chart recording.

All square specimens

IS
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16
Table 7. Compressive Properties of Composites (Maxima of Ultimate Strength Values)

Temperature

(K)

Ultimate Strength, a^^

(108 N/n,2) (io3 psi)

Ultimate

Compression,

(%)

Number of

Specimens

295

76

4

295

76

4

295

76

4

295

76

4

295

76

4

295

76

4

295

76

4

295

76

4

5.6 mil Boron-6061 Aluminum

28.7

32.6

2.60

4.56

6.56

36.6

2.27

4.56

4.52

5.62

16.5

15.4

1.05

3.11

2.99

Longitudinal (0°)

416

473

Transverse (90°)

37.7

66.2

95.1

0.5

0.7

0.5

0.9

1.2

5.6 mil Boron-5505 Epoxy

Longitudinal (0°)

531 0.6

Transverse (90°)

32.3

66.2

65.5

S Glass-NASA Resin 2

1.6

1.5

1.1

Longitudinal (0°)

81.6

239

224

Transverse (90 )

15.3

45.1

43.4

1.3

3.0

2.6

3.4

1.1

1.1

Type AS Graphite-NASA Resin 2

Longitudinal (0°)

5.61

12.3

8.83

1.04

1.51

1.50

81.2

178

128

Transverse (90°)

15.1

21.9

21.8

0.5

1.0

0.8

0.8

1.1

1.0

17
Appendix I: Calculation of Compressive Moduli

The elastic (Young's) modulus, E, is determined from a stress-strain

curve by the relationship:

Ae

where L = load, A = specimen cross-sectional area and e = specimen strain.

If the only available record of the test is a strip chart trace of load as

a function of time, the time axis must be calibrated in units of strain.

For an infinitely rigid test machine, the strain would be given by the

product of time and crosshead speed divided by the specimen gage length. In

reality, a fraction of the crosshead motion causes strain in various components of the load train other than the specimen. This extraneous strain may

be accounted for if one knows the functional relationship between incremental

specimen strain, Ae, and incremental strip-chart time. At:

Ae = f(At) (2)

The factor f is a function of both the spring constant of the loading

system and of specimen stiffness, E'A. It is uniquely defined for a given

specimen and load train, and may be determined by comparing a strip-chart

trace with a load-strain trace recorded concomitantly from a strain-gaged

specimen. Specimens of differing stiffness yield differing f factors,

reflecting differences in the distribution of overall strain between the

specimen and the load train. However, if f factors are calculated from a large number of specimens which vary in modulus and cross-sectional area

(but are of identical gage length) , a plot of f versus the E"A product

provides a calibration curve for the test system.

Combining Eqs. (1) and (2) yields:

^ A • f(At)

Since AL, At and A are known, an iterative process is used to find the

correct value of the factor f. An estimate is made of the modulus, the

E'A product is calculated, and the corresponding f value is determined from

the calibration chart. This value is substituted in Eq. (3). A few iterations suffice to narrow estimated and calculated moduli to within + 10%,

which appears to be a reasonable limit of accuracy for this method.

The same f value may be used to estimate the fracture strain from the

time scale of a load-time trace.

It is also possible to estimate the ultimate strain from the total test

time using Eq. (2) . In those cases where a strain gage was applied but failed

before specimen fracture, a comparison of the initial linear portion of the

load time curve with the modulus from the load-strain curve yields an f value

from Eq. (3) without recourse to the calibration curve or any iterations.

18
Reinforced Structural Composites at Cryogenic Temperatures -2: _ Advanced Composites/' Cryogenics 15(12), pp. 701-722 (1,975).

The low-temperature mechanical and thermal properties of advanced-fibre reinforced

structural composites are reviewed. The magnitude and range of particular properties

are discussed with respect to composite type and temperature. A property—material

cross reference is given with a 128-entry bibliography. This is Part 2 of a two-part

series. Part 1 considered glass-reinforced composites.

Mechanical and thermal properties of filamentaryreinforced structural composites at cryogenic

temperatures

2: Advanced composites

M. B. Kasen

Nomenclatu re

0tU tensile ultimate strength

^1 imti^il tpncilp mrvnlililc E\ secondary tensUe modulus

tensile ultimate strain

flexural ultimate strength

initial flexural modulus

secondary flexural modulus

compressive ultimate strength

compressive modulus

compressive ultimate strain a" interlaminar shear strength

fatigue strength

(jby bearing yield strength

bearing ultiinate strength

a' impact strength

X thermal conductivity

MIL tliermal contraction (expansion)

Cp specific heat at constant pressure

P density

This review has four main objectives: (a) to give the designer some idea of the general magnitude of property values which

can reasonably be expected from a given category and class of advanced-fibre composites within the cryogenic range;

fb) to provide insiglit into the ranking of specific composite

classes with regard to a specific property; and (c) to allow

liim to assess whether the property of interest is likely to

increase, remain unaffected, or decrease with lowering of

temperature. 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 1 and

The author is in the Cryogenics Division, NBS Institute for Basic Standards, Boulder, Colorado 80302, USA. Received 16 June 1975.

includes references to both the glass and advanced- fibre reinforced composite literature. The scope of the literature survey has already been described in Part 1 (CRYOGENICS

Vol 15, No 6 (June 1975) pp 327-49)

We define a composite category by the general reinforcement

type, for example, glass-fibre or advanced-fib re (graphite,

boron, etc). We subdivide the category into composite

classes by the general matrix type, for example, glass- polyester or graphite-epoxy. We subdivide the class by re- ferring to a composite type when a specific reinforcement/

matrix combination is specified, for example, HT-S/X-904

epoxy. The term 'advanced-fibre' is used to distinguish fibres having

high modulus (20-70 x 10^ lb in'^, 138-483 GPa) in contrast to the relatively low modulus of glass fibre (10 x 10^

lb in'^, 69 GPa). Fibres of boron, graphite or proprietary

organic type dominate the advanced-fibre field.

Several differences will be noted between the presentation of

the data in this work and in Part 1. First, because of the wide

variation in properties among the advanced fibres (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 advancedfibre composites at cryogenic temperatures, present data are

restricted to uniaxial longitudinal and uniaxial transverse layups. Thirdly, the overwhelming majority of the relevant data for advanced-fibre polymeric systems are reported for epoxy matrices, in contrast to the variety of polymeric matrices for which data were available in glass-reinforced systems.

Finally, data are presented for advanced-fibre reinforcement

of a metal matrix (aluminium).

A more subtle difference in the present work is the separ- ation of static mechanical properties into primary and secon- dary categories. The former category distinguishes those

properties useful in design calculations, while the latter dis- cusses properties more qualitative in nature, which are pri- marily useful in quality control or for comparative perform- ance screening of composites.

In other respects, the data presentation follows that used in Part 1, that is, hterature 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.

CRYOGENICS . DECEMBER 1975 19 701
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 complexity of the subject and the necessity of a cut-off data for data acquisition 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 this review reflect the published results of

the cited authors. These data have not been experimentally

verified by NBS, and the conclusions and evaluations present- ed herein do not imply approval, endorsement, or recommendation of any commercial product by NBS.

There exists no universally accepted method of determining

many of the reported properties. It is not easy to obtain

valid uniaxial longitudinal tensile fracture because of the

difficulty of transferring the tensile load from the specimen

grips into the fibres of the specimen without introducing

excessive localized stress concentrations. Problems also arise , in uniaxial longitudinal compression testing where an un- supported specimen may fail by the '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.

The discussion of properties of the composites included in

this review does not take into consideration the effect of variations in fibre resin ratio among specific types of composites and test specimens. Boron-reinforced composites

normally contain an 0.45—0.50 fibre resin ratio volume, while composites reinforced with graphite or organic fibre conventionally have an 0,50—0.65 ratio. Composite properties are strongly influenced by this ratio. Controlled variations in many of the properties are obtainable in practice by specific variation of the fibre content of the composites. The property

data discussed in this paper reflect values and trends reported

for composites containing of the order of 50% fibre volume. However, the reader must refer to the literature for specific

composite data.

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. Recent introduction of new reinforcing fibres, improvements

in composite fabrication techniques and refinement of test methods have resulted 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 per- spective on the field. The two primary reinforcement sy- stems are boron filaments produced by vapour deposition of boron on a very fine tungsten wire substrate and graphite

^ibre produced by graphitization (pyrolysis) of an organic precursor fibre. Additionally, a proprietary aromatic polyamide organic fibre commercially called Kevlar 49 (PRD 49)

is currently receiving attention. The boron filaments were

initially produced at 4 mil (1 mil = 25.4 /xm) diameter; however, 5.6 mil diameter filaments are now widely used. Boron

fibre coated with a thin layer of silicon carbide, produced

under the trade name of Borsic, is reported to have improved

interfacial bonding to certain matrix types. The graphite

fibre is much more complex, as it is not only possible to

produce such fibres from different precursor materials (the two most common being rayon and polyacrylonitrUe (PAN)),

but it is also possible to vary the production process to pro- duce fibres 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 fibres has proceeded at a very rapid rate. Today the user is confronted with an abundance of fibre choices, many of which are not well characterized and many of which will disappear to be replac- ed by newer types. At the present time, the potential user of

these materials would be wise to restrict his interest to those

types whose behaviour is reasonably well known, barring

compelling reasons for doing otherwise. Graphite fibres may be produced with elastic moduli varying

from 25-75 x 10^ lb in^ (170-500 GPa), with strengths

varying inversely to moduli. In the present report, these

fibres are classified according to low, medium or high modulus, the differentiation being < 40, 40-60 and > 60 x 10^

lb in-2 , respectively (< 275, 275-41 5, and > 4 1 5 GPa). 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.

Primary static mechanical properties

The key static mechanical properties required for preliminary

design calculations with composites are: uniaxial longitudinal

and transverse tensile and compressive ultimate strengths; uniaxial in-plane shear ultimate strength; and ultimate strains

in the uniaxial longitudinal and transverse directions. Key

elastic properties are uniaxial tensile and compressive moduli

in the longitudinal and transverse direction, uniaxial in-plane shear moduli and Poisson's ratios in the longitudinal and

transverse uniaxial direction (see reference 124, Vol 1, Design). The literature data available on these key properties for composites within the cryogenic range are far from

complete, and in some cases, non-existent. Nevertheless,

available data provide a feel for the magnitude and temperature dependence of many key properties and serve to define , areas in which further data are required. For most mechanical properties, the data for the graphite- reinforced composites have been separated from that of the other advanced composites, that is, those reinforced with

boron or PRD 49 (Kevlar 49). This reflects the larger amount of available data on the graphite-reinforced materials

Table 1. Typical composite densities

Composite

system

Fibre/resin

ratio

Density, lb in'^

(g cm"3)

s-glass-epoxy 0.60—0.67

Kevlar 49-epoxy 0.60-0.65

Boron-epoxy 0.55

Graphite-epoxy 0.55—0.60

Boron-aluminum 0.50

(4 mil)

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)

702 20 CRYOGENICS . DECEMBER 1975 i
170

160-

150-

140-

130-

120 MO

100 90

70-

50-

—— High modulus —— Medium modulus Low modulus

* Minimal doto

X 2a

2b

2c

xlO 6

50 100 150 200

Temperolure, K

250 300

Fig.1 Ultimate tensile strength, a*", of graphite-epoxy

composites

a - Uniaxial lonqitudinal: 1 - Thornel 75 ^'; 2 - GY-70

*"';4 - HMG - 50 '°;5 - MH-S '°;4^

3 - Thornel 50 6 - Modmor I * : 7 - Samco 320 ° : 8 - Thornel 25 '

, 9 - HT-S 9,10,13,13.2,58. -,0 - Modmor II ^ 11 - HMG-25

b - Uniaxial transverse: 1 - Thornel 75/ERLB 461 7 ; 2a - HT-S/NASA Resin 2

'°; 2b - HT-S/4617 (modified) 2c - HT-S/X-904

and the large variety of graphite fibres for which data have been pubUshed.

Composite tensile strength

Fig.l summarizes the available data on the temperature depen- dence of the ultimate tensUe strength of graphite-reinforced epoxy composites from room temperature into the cryogenic

temperature range for the uniaxial longitudinal and uniaxial

transverse orientations. A surprisingly large amount of cryogenic data were found in the literature for graphite-reinforced

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 fibres. The available data summarized on Fig.l indicate that graphite- reinforced epoxy composites may suffer significant strength

losses upon cooling to 77 K, and as was observed with the

glass-reinforced materials in Part 1 , the strength behaviour

below 77 K appears to become erratic.

Fig. la indicates that the uniaxial tensile strength obtainable m graphite-reinforced composites is about 30% of that obtainable with glass-fibre reinforcement in the cryogenic range. The ultimate tensile strength of composites tested in the uniaxial longitudinal mode should be fibre controlled; hence,

the tensile strengths should inversely follow the modulus of

the fibre. This trend is not reflected in the data of Fig. la -

indeed the averaged data indicate the converse; higher

strengths are associated with higher modulus fibres and

lower strengths with lower modulus fibres. 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 strength values reported for each modulus range. Thus, the higliest overall ultimate strength

at 77 K was 147 x lOMb in'^ (1.01 GPa) reported for the low modulus HT-S graphite fibre in X-904 resin. The

next highest was 130 x 10^ lb in'^ (0.90 GPa) reported for HMG-50 in a flexibilized epoxy matrix '° for Modmor I and

for Samco 360 in modified ERL 1 1 56,^ all medium modulus

fibres, while 126 x 10^ lb in"^ (0.87 GPa) was the maximum

reported for the high-modulus Thornel 75 fibre in ERLB

4617.5' This suggests that the higher strengths normally

associated with the lower modulus fibres may have been

lost in the averaging process due to variations in composite

quality or perhaps due to difficulties in obtaining valid tensile fractures. Additional evidence that the test method is capable of affecting the results is found in the work of

Larsen and Simon who report almost diametrically

opposite temperature dependence of the ultimate tensile strength for NASA Resin 2 and ERLB 4617 epoxy rein- forced with HT-S fibre when each was tested, first as flat tensUe coupons and subsequently as NOL (Naval Ordnance

Laboratory) ring specimens. NASA 2 resin is a bisphenol

A epoxy system modified for low temperature flexibility by means of a long-chain anhydride and a high molecular weight

tricarboxy acid. This resin, consisting of Epon 828/DSA/

Empol 1040/BDMAin proportions 100/115.9/20/1 pbw,

was developed by Soffer and Molho ^ under NASA

Sponsorship.

The uniaxial longitudinal tensile strength data for other types of advanced composites are presented in Fig.2a. The data

for PRD 49-epoxy, boron-epoxy, boron (Borsic)-aluminium,

stainless steel-aluminium, and the hybrid Borsic-steel- aluminium and Borsic-titanium-aluminium composites all have higlier absolute values of tensile strength and retain

their strength to lower temperatures than do the graphite- reinforced composites. Data for Borsic and boron-fibre composites have been combined except where differences

in reported values justified separating the data. The hybrid

Borsic-steel-aluminium (6061) composite contained 4.2 mU

Borsic fibres in the tensUe direction and stainless wires in the

transverse direction. The hybrid Borsic-titanium-aluminium (6061) composite contained 4.2 mil Borsic fibres in the tensile direction with j3-3Ti foil interleaved between the boron

lamellae.'^ While the latter two composites are not strictly uniaxial, they have been included to illustrate the interesting

possibilities of hybrids.

The magnitude and temperature dependence of the ultimate

tensile strength of the graphite HT-S/polyimide (Skybond

703) material were comparable to that of the graphiteepoxies. A very sharp rise in strength on cooling was re- ported for the HT-S/epoxy-phenolic (HT424 Primer).

The stainless steel reinforced composite contained NS-355

stainless wires in a 2024 aluminium 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 aluminium, which has

similar strength properties.

Boron-aluminium composites are available commercially with

CRYOGENICS . DECEMBER 1975 21 703 !
0 50 100 150 200 250 300

Temperoture , K

Fig. 2 Ultimate tensile strength, a , of miscellaneous advanced

composites

a - Uniaxial longitudinal: 1 - PRD 49-1 /epoxy

o u 13 22 47 53 58 81 T lit o / i j 13 2 — boron-epoxy . . / °. ; 3 — HT-S/polymide ; 4 — HT-S/epoxy-phenolic 5 — boron-aluminium (6061

)

6 — Borsic-steel-aluminium (6061)'^; 7 — Borsic-titanium- aluminium (6061 ) ; 8 — stainless steel-aluminium (2024)^.^"* b — Uniaxial transverse: 1 - boron-epoxy ; 2 - boron- aluminium (6061 ) 3 — Borsic-titanium-aluminium

either 4 or 5.6 mil boron or Borsic reinforcement, while 8 mil boron fibers are currently being studied. Available liter- ature values of uniaxial tensile strength at 77 K ranged from

163-202 X 103 lb in-2 (1.12-1.39 GPa), with highest values reported for 5.6 mil Borsic/6061.'3 Data were not available

for any other aluminium alloy matrix. Boron-epoxy composites containing either 4 mil or 5.6 mil fibre are commercially available in the form of prepreg tape, that is, with the

plastic matrix partially cured to facilitate component fabric- ation and to improve composite quality. The literature values

for boron-epoxy composites tested in the longitudinal dir- ection at 77 K varied from 167-226 x 10^ lb in-2(1.15-1.56

GPa), the highest value being reported for the commercial

SP-272 product.'

3

PRD 49 (Kevlar 49) is a relatively new type of organic fibre. Present data indicate that this fibre in an epoxy matrix is capable of developing tensile strengths at cryogenic temperatures comparable to that developed with the graphite

fibres. Hoggatt reports NOL ring ultimate strengths of 146 x 10Mb in-2 (1.0 GPa) for PRD 49-1 in a NASA Resin 2 matrix and 183 x 10^ lb in'^ (1.26 GPa) 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-aluminium and Borsic-titanium- aluminium complete the group of advanced composites that has reasonably high strength in uniaxial tension. The some- what lower strength of these hybrids, compared to the conventional boron-aluminium composite, reflects the reduced

density of boron fibres in the loading direction. Metal-matrix composites are clearly superior to polymermatrix composites in transverse mechanical properties. As

seen in Fig.2(b), the ultimate uniaxial transverse tensile strength of boron/6061 aluminium is about 50% higher than

that of boron-epoxy. Addition of titanium foil to the boron- aluminium further increases the transverse 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 conventional boron-epoxy composite is much lower than that of the metal-matrix materials, values

are still in excess of those developed with the graphite fibre reinforcement. The boron/6061 aluminium data reviewed here reflects the

as-fabricated (F) condition. Solution-treating and aging the composite increase the transverse ultimate strength by about 30%, with a slight decrease in transverse ductility.

The literature reported a 77 K transverse ultimate strength of 14.2 X lOMb in-2 (97 MPa) for 4.2 mil Borsic/6061 alumin- ium and a somewhat higher value of 24.9 x 10^ lb in"^ (172

MPa) for 5.6 mil boron reinforcement. It is also probable

that this strength difference is real; however, available data 13'126 jIjq indicate that the observed difference may

be due to residual stresses inherent in the small diameter

fibre rather than to the presence or absence of a silicon car- bide 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 fibre reinforcement. In contrast to a maximum of about 10 x 10'

lb in'^ (69 GPa) for glass-reinforced composites in the fibre direction. Figs 3a and 4a show that modufi ranging from

30-40 x 10^ lb in-2 (207-276 GPa) are obtainable with

several of the advanced fibre composites. Thus, while glass- reinforced composites may at their best equal the modulus

of aluminium, the modulus of composites using advanced

fibres may at their best equal and often exceed that of steel. Looking first at the graphite-epoxy composites, Fig.3a shows

that the uniaxial longitudinal modulus of the composite does, on the average, reflect the modulus of the fibre when tested in the uniaxial longitudinal mode. The available data indicate

that the modulus of the Thornel 75 composite (the only re- presentative of the high modulus fibers) declines to about

that of the medium modulus fibres at 77 K. With this exception, the average data indicate that the moduli in the fibre direction slightly increase on cooling to 77 K. No correlation

is observed between fibre modulus and transverse composite

moduh in Fig.3b; indeed, none is expected, as the latter is controlled by the properties of the matrix and by the fibre- matrix interfacial bond strength.

A comparison of Figs 3 a and 4a shows that the uniaxial longi- tudinal modulus of graphite-epoxy composites may exceed that of any other type of advanced composite when certain graph704 22 CRYOGENICS. DECEMBER 1975

4 %
xlO50 F"

45

40

35

30

25

5 20

^ 15

a

'° xlO^2

^— High modulus — — Medium modulus Low modulus •

* Minimol doto

xlO" 35

30

25

20

10 ^

1 .

1 1

2a

2b

50 100 150 200 250 300

xlO

12

Temperature, K

Fig. 3 Initial tensile modulus, E\, of graphite-epoxy composites

a — Uniaxial longitudinal: 1 — Thornel 75 ^'

; 2 — Thornel 50 ; 3 - Modmor I ';4 - Sancom 320 **; 5- HT-S ^-'''.'^.-e- HMG-25 ** b - Uniaxial transverse: 1 - Thornel 75/ERLB 461 7

^'

;

2a - HT-S/NASA Resin 2 '°;2b - HT-S/4617 (Modified); 2c - HT-S/X-904

ite fibres are used. However, a comparison of Figs 3b and

4b shows that the uniaxial transverse tensile moduli of graphite-epoxy composites are much lower than that develop- ed by the other advanced composites over the entire cryogenic

temperature range. A closer look at the literature data from which Fig.3a was

prepared revealed that among the medium modulus fibres, the modulus at 77 K was 39-43 x 10^ lb in'^ 269-297

GPa) for Samco 360; the highest value being reported with

a modified ERL 2256 resin. ^ The Modmor I fibre produced

almost as high moduli, ranging from 32—42 x 10^ lb in'^ (221-290 GPa), with the highest values reported for a NASA Resin 2 composite.^ (Note that these same fibres developed high uniaxial tensile strengths). A distinctly lower range of 26-33 x 10^ lb in^ (179-228 GPa) was

reported for Thornel 50; the highest values were for a modified ERL 2256 composite.^ among the low modulus

fibres, HT-S (the most tested fibre) produced 15-20 x 10^

lb in"2 (103-138 GPa), the highest value being reported with X-904 resin,! 3 while 17-22 x 10Mb in'^ (1 17-152

GPa) was the reported range for HMG-25, the highest values being obtained in a NASA Resin 2 matrix.^ As the transverse modulus properties are matrix dominated,

Fig.3b shows significant differences for the same HT-S fibre

in different epoxies. The Ulustrated data are for the same

composites as appear in Fig. lb. Here the lower moduli of

the HT-S/NASA Resin 2 and the HT-S/4617 (modified) are due to the additions of flexibilizers 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 composite modulus value was reported for the high-modulus

Thornel 75 fibre in ERLB 4617 resin than for the HT-S/X-904 combination.

Leaving the graphite-reinforced composites and turning to those reinforced with other advanced fibres, we observe on

Fig.4a that the uniaxial longitudinal moduli reported for the boron-epoxy, the boron-aluminium and the hybrid

Borsic-titanium-aluminium and Borsic-steel-aluminium composites all cluster around 30-35 x 10^ lb in'^ (207—242

GPa) with little temperature dependence. The PRD 49/

epoxy modulus is reported to be significantly lower, at 16-18 X 10^ lb in-2 (1 10-124 GPa), but undergoing a

significant rise between 77 K and 20 K. In general, these

uniaxial longitudinal moduli values compare favourable with those of the graphite-epoxies.

The uniaxial transverse moduli are another matter. As seen

in Fig.4b, the transverse moduli of the aluminiurn-matrix

materials are far higher than those of the graphite-epoxies

and are, in fact, almost twice as high as that of unreinforced 6061 aluminium. The transverse moduli of the boron-epoxy

is much lower than that for the same fibre in an aluminium

matrix, but is still almost twice that of the graphite-epoxy

materials. These data reflect the contribution made by

fibre-matrix interface to the transverse modulus.

100 150 200

Temperature , K

Fig.4 Initial tensile modulus, of miscellaneous advanced

composites • a — Uniaxial longitudinal: 1 — boron-epoxy ; 2 — PRD 49—

1 /epoxy 3 — boron-aluminium (6061 ; 4 — Borsic- titanium-aluminium (6061)'''; 5— Borsic-steel-aluminium (6061

)

13

Uniaxial transverse: 1 — boron-epoxy (SP-272, Narmco 5505) 81 2 — boron-aluminium (6061 ) ; 3 — Borsic-titanium-aluminium (6061)'';4 - Borsic-steel-aluminium (6061)'^

CRYOGENICS. DECEMBER 1975 23 705
~ 200-

50 100 150 200

Temperoture, K

250 • 300

Fig. 5 Ultimate compressive strength, a"*", of advanced

composites

13 SI a — Uniaxial longitudinal: 1 — HT-S/X-904 ; 2 — boron-epoxy

(SP-272, Narmco 5505); 3 - boron-aluminium (6061); 4 — Borsic-titanium-aluminium (6061) ; 5 — Borsic-steel- aluminium (6061 )

' b - Uniaxial longitudinal: 1 — HT-S/X-904 ; 2 — boron-epoxy , (Narmco 5505, SP-272) ^'

; 3 - boron-aluminium '

Looking in more detail at the uniaxial longitudinal data, we

find that the boron (Borsic)-aluminium data ranged from

30.4-36.2 X 10^ lb in'^ (210-250 GPa) with Httle temperature dependence. These data combine values obtained from fibres of 4 mil and 5.6 mil diameter in the two production variants, as no significant difference was reported for these materials. The boron-epoxy data represent average

results reported by Nadler et al ^' on SP-272 and Narmco

5505 (data range was not available). The hybrids, Borsic- titanium-aluminium, and Borsic-steel-aluminium, were the same composites discussed in reference to the tensile data of Fig.2a. The slightly lower uniaxial tensile modulus of

the hybrids compared to that of the conventional boron- aluminium reflects the lower density of fibres in the stress direction in the hybrids.

The PRD 49 data reflect input from both an ERLB 4617

and a NASA Resin 2 matrix.59.60 xhe ERLB 4617 composite produced a slightly increased tensUe modulus at all temperatures, its value being 18.5 x 10^ lb in'^ (128 GPa)

at 77 K compared to 17.5 x 10Mb in^ (121 GPa) for the NASA Resin 2. A large rise in modulus between 77 K and

20 K was reported for both matrices. Considering the transverse data of Fig.4b in more detail, we observe that the three aluminium-matrix composites

have similar moduli, about 20 x 10^ lb in"''- (138 GPa).

The boron-aluminium 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 Fig.4b again reflect the average values obtained

from the commercial SP-272 and Narmco 5505 materials.^'

Composite 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 Figs 5 and 6, combining the graphite-epoxy data with those of the other composite

types. These data show that the advanced composites have signi- ficantly higher compressive strengths and moduli than those of the glass-reinforced composites. Comparison of Fig.5 with Figs 1 and 2 also shows that the compressive strengths

are 100-200% higher than the tensile strengths for the same

composites.

The higliest uniaxial longitudinal compressive strength, more than 200% above that obtainable with glass reinforce- ment, was reported for the commercial boron-epoxy types

SP-272 and Narmco 5505. The compressive strength of

this composite class was observed to increase on cooling to 77 K. Boron-aluminium reportedly has a comparatively lower uniaxial compressive strength. The latter data reflect input from both 4.2 mil Borsic and 5.6 mil boron-fibre composites.'^ A slightly higher strength was reported for the 5.6 mil composite, but not sufficiently to justify a sep- arate plot. Again, the Borsic-titanium-aluminium and

Borsic-steel-aluminium hybrids developed strengths somea isi50 100 iSc 200

Temperature , K

250 300

Fig. 6 Compressive modulus, f*^ of advanced composites

a - Uniaxial longitudinal: 1 - HT-S/X-904 '^; 2 - borOn-epoxy

(Narmco 5505 SP-272)*' ; 3 - boron-aluminium (6061

)

4 — Borsic-titanium-aluminium (6061 ) ; 5 — Borsic-steel- aluminium (6061 )

' b - Uniaxial transverse: 1 - HT-S/X 904 '

; 2 - boron epoxy

(Narmco 5505, SP-272)'" ; 3 - boron-aluminium (6061)

706

24

CRYOGENICS,. DECEMBER 1975
what below that of the conventional boron-aluminium, re- flecting the decreased density of boron fibres in the stress direction. Lowest uiuaxial longitudinal compressive strengths were reported for the HT-S X-904 graphite-epoxy composite;

however, this composite developed a compressive strength

about twice that developed in tension and about 405c higher than that available with glass reinforcement. The compressive strength of the graphite-epoxy composite was reported

to increase v.ith cooling, in contrast to the tensile behaviour. The available data on the transverse compressive strengths of boron-epoxy, boron-aluminium, and graphite-epoxy composites are compared on Fig.5b. These composites appear

in the same relative order of strength as in the longitudinal

test mode. The commercial SP-272 and Narmco 5505

boron-epoxy products display substantial increases in compressive strength on cooling. Lesser increases are observed

in the boron; aluminium and in the HT-S X-904 graphite- epoxy composite.

The compressive modulus of a given composite is e.xpected

to be the same as the tensile modulus. This is found to be

generally true for the HT-S X-904 graphite-reinforced material, for the boron-epoxy, for the Borsic-aluminium composites, and for the Borsic-steel-aluminium hybrid, as may be

seen by comparing the compressive moduli of Fi2.6a v,ith the tensile moduli of the same composites on Fig 3 and 4. The Borsic-titanium-aluminium data are at variance with

this principle, as the reported compressive modulus is about

305t higher than in tension. The boron-aluminium and the boron-aluminium hybrids again show the highest longitudinal

moduli, with the boron-epoxy only slightly lower. The

modulus of the HT-S X-904 graphite-reinforced 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-aluminium is found to have a higher transverse compressive modulus than either the boron-epoxy or the

graphite-epo.xy composite.

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 reflects the work of Hertz et

al.'3 Nevertheless, these data suffice to illustrate that: (a) the strain at tensile fracture is ver\' low for the advanced

composites, being of the order of 1—9 x lO''^ in the longi- tudinal uniaxial direction; and (b) the strain at tensile frac- ture is relatively independent of temperature within the

cry ogenic range. These data are summarized on Fig.7. These

strains are about an order of magnitude lower than those obtained with glass-reinforced composites, retlecting the

negligible fracture strain of the advanced fibres.

Fig.7 includes data on one graphite-reinforced composite

I HT-S X-904). on two conventional boron-aluminium composites (4.2 mil Borsic and 5.6 mil boron, both in 6061

aluminium) and on the two Borsic-aluminium hybrids dis- cussed previously with reference to the other properties.

Both longitudinal and transverse ultimate fracture strains appear on Fig.7.

Fig." indicates that the transverse fracture strain of the HT-S'X-904 graphite-epo.xy composite wiU be about 2 x 10^.

or about 259c of the longitudinal strain. By comparison, the

transverse fracture strain for the commercial HT-S'30O2

5C

Fig.7 Ultimate tensile strain, (uniaxial tensile)

13

, of advanced compKDSites

,13 5.6 mil boron-aluminium (6061 4 - 4.2 mil Borsic-

)

titanium-aluminium (6061) 5 — 4.2 mil Borsic-steel- aluminium (6061)

graphite-epoxy composite is reported to be 7.8 x 10"^, only slightly less than its longitudinal fracture strain of 8.6 X lO"-'. The interfacial-bond strength between the fibre and

matrix in the FTT-S 3002 composite was undoubtedly superior to that e.xisting in the HT-S 'X-904 composite.

Fig.7 shows that a 5.6 nul boron/606 1 aluminium composite is expected to fail at about the same strain in both

the longitudinal and transverse directions. The 4.2 mil

Borsic; 6061 composite is e.xpected to fail at a much lower

strain when tested in the transverse direction. Hertz et al found that the 4.2 mil Borsic fibre was failing by longitudinal fibre spUtting during transverse testing, while such spUtting did not occur in the larger diameter fibre. The authors concluded that the problem was related to the use of plasmasprayed tape. Higher transverse strains were observed in the same type of composite when diffusion-bonded tape was

used to make the test specim.ens. They concluded that diffusion-bonded boron-aluminium tape was superior for oriented-ply appUcations. Kreider and Prewo ^ - ' has pro- posed an alternative explanation for longitudinal spUtting of the small diameter boron fibres, citing results of diametral- compression tests of individual fibres, which indicate that such splitting reflects strength anisotropy in the sm.all fi.bres due to residual stresses retained from the original fibre manufacturing process and from pre-existing flaws in the

fibres. Such anisotropy was not observed in the larger fibres. Fibre spUtting was not directly related to the presence or absence of a siUcon carbide coating on the fibres. (Manufacturers of boron-aluminium tape report that current fab- rication technology has overcome the problem of fibre

splitting)

The Borsic-titanium-aluminium hybrid reponedly developed

higher strain values in both directions than did the conventional 4.2 mil Borsic-aluminium, but lower values than that of the composite reinforced with 5.6 mil boron. The borsicCRYOGENICS . DECEMBER 1975 25 707
20

18

16

14

12

10

8

6 -

—— Transverse

Longitudinol

50 100 150 200 250 300

Temperolure, K

Fig.8 Ultimate compressive strain, e^, of gra^hite-epoxy

composites: unaxial compression HT-S/X 904

steel-aluminium hybrid shows clearly the increase in trans- verse 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. The only data available on the temperature dependence of compressive strain within the cryogenic temperature range was for the HT-S/X-904 graphite-epoxy composite, again

reflecting the work of Hertz et al.'^ These data, Fig.8, 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. As in the tensile case, these strain values are substantially

lower than those reported for an HT/S/3002 graphite- reinforced composite.'^'*

Composite in-plane shear strength and modulus

The key shear properties are the in-plane ultimate strength

and modulus, sometimes referred to as the intralaminar or longitudinal shear properties. The requirements for obtaining valid shear property data are the same as for metals, that

is, values must reflect pure shear resolved onto a 45° shear plane with respect to the tensile axis without significant

compressive or tensOe forces normal to the shear plane. Three shears may be defined in a bulk, uniaxial composite

laminate. Most important are the two longitudinal shears represented by shear between parallel fibres in the plane of

the fibres. For example, taking the x axis parallel to the

fibres, longitudinal shear strengths may be defined as a*^ or o*^ ; however, in thin laminates, only one such in- plane shear is of consequence. A third shear, which is trans- verse or cross-fibre, is not an important composite design

parameter.

It is important to distinguish between in-plane or intrala- minar shear defined as above and interlamimr shear, which

refers to shear between adjacent layers in a layered composite. The latter, sometimes called horizontal shear, is not a basic material property and is used primarily as an

adjunct to quality control or in material screening tests. Because the test is simple and inexpensive, a large amount

of interlaminar shear data have been reported.

In-plane shear is a difficult property to measure accurately because of the necessity of avoiding extraneous compressive

or tensile stresses on the shear plane. The torque tube

(torsion) method works well but is expensive. Consequently,

somewhat simpler methods such as the 'rail shear' or the

'+ 45° laminate' test have been devised. The interested reader can find discussions of the relative merits of these

latter methods in ASTM STP Vols 460, 497, and 546. Few data are available on the in-plane shear properties of composites at cryogenic temperatures. Hertz et al '-^ report

a decrease in the shear strength of HT-S/X-904 graphiteepoxy from 8 592 lb in'^ (60 MPa) at 295 K to 7 378

lb in'2 (51 MPa) at 77 k, using the torsion tube test method.

A concomitant shear modulus increase from 7.27 x 10^

Ibin"^ (10.5 GPa) was reported. Obviously, much additional

in-plane shear data are required before trends in the cryogenic temperature dependence of shear properties can be

defined. Room temperature values of shear strengths and

moduli are reportedly about 18 x 10^ lb in'^ (124 MPa)

and 9.5 x 10^ lb in"^ (65 GPa), respectively, for boron- aluminium, while values for boron-epoxy are about 15x10^

lb in-2 (124 MPa) and 7 x 10^ lb in'^ (0.48 GPa), respectively.'^'' Reported room temperature values for the shear

strength of graphite-epoxy composites vary from 9— 13 x

103 lb in-2 (62 MPa) with a modulus of about 6.5 x 10^

lb in"2 (0.45 GPa), the variation in strength reflecting variation

in properties of the graphite fibre used in the composite. As most of the other mechanical properties of boron-aluminium

and boron-epoxy composites either remain unaffected or increase upon cooling to cryogenic temperatures, these room temperature values likely represent conservative values

for such materials at the lower temperatures. However,

confimiation must await further testing at the cryogenic

temperatures.

Secondary static mechanical properties

The properties discussed in this section are classified as se- condary only in that they are not generally useful in a pre- dictive design analysis using macromechanics composite

theory. However, such secondary properties are important

in other ways. Because composites are largely fabricated in

situ, a large number of options are available among fibre (or combinations of fibres) and matrix materials. Options are

also available in fabrication procedure. It is necessary to

efficiently evaluate various combinations in a meaningful

comparative manner in order to optimize the final selection. Furthermore, having made a selection, it is necessary to have viable techniques for measuring the consistency of composite quality over a production run. For such objectives, the flexure test and the interlaminar shear test (usually also performed in flexure) are relatively inexpensive and informative. Additionally, there are properties that

are classed as secondary only in that the results are specific to a given geometry as, for example, bearing yield strength.

Composite flexural strength and modulus

The flexure test requires that a bar or plate specimen be supported near its ends while deflected by a moving ram at

its cenre, providing symmetrical three-point loading. Dimensions of the fixture and specimen are proportioned to deve- lop maximum stress in the outer (tension) fibres as flexure

progresses. Fibres run lengthwise along the specimen, that

is, between the supports, in uniaxial longitudinal flexure and

across the specimen in uniaxial transverse flexure. Fewer data were available on the temperature dependence of

708 26 CRYOGENICS . DECEMBER 1975
(4

MedtUTT moOiius

5CFig.g Ultimate flexural strength,

a — Uniaxial longitudinal: 1 — GY-70

^raphlte-epoxy composites

; 2 - Thornel 50

3 - Courtaulds HM ^,-4 - Fibralloy 300 '^;5 - HT-S 6- Modmor II '^;7 - HMG-25

b - Uniaxial transverse: 1 - Modmor 11/1004 *^;2 - HT-S/1004

3 - HT-S BSP-2401 '^;4 - HT-S/X-904 5 - HT-S/3002

6 - GY-70, 1004 '^;7 -GY-70/E-350A ''';8 - GY-70 X-904 " -'^

the llexural strength and elastic modulus than for tensile properties; furthermore, most of those available were re- ported for graphite-epoxy composites. Fig.9 summarizes

the flexural strength data for the latter, while Fig. 10 summarizes the available data for other t\ pes of reinforce- ments or matrices. The minimal available data on flexural modulus properties appear in Fig. 11. Comparison of Fig.9 with Fig.l shows that the graphite- epoxy flexural strength data span a much greater range of

values than does the tensile strength data in both the longi- tudinal and transverse directions. The e.xpected higher

strength of the lower modulus fibres is more in evidence

for the flexural test than it was for the tensile test mode.

B\' far the largest amount of data was available for HT-S '

epo.xy composites. Furthermore, such composites had the

highest longitudinal flexural strength of all the graphiie- reintorced materials for which data were available. Strengths were reported over a 113-234 x lO^ lb in'^ (0.78-1.61

GPa) range at 77 K. This is significantly h fter than the 325-^70 X 103 lb in-2 (2.24-3.24 GPa) range reported

for unia-xial glass-epo.xies in Part 1 . These data included

lest series designed to investigate cure cycles, aging effects, and environmental effects. ^ ^ 3-- ^ Most of the work was

done with E-350, X-915, or X-904 matrices, for which the

average flexural strengths at 77 K were reported to be very

simUar. varving only from 167—178 x 10^ lb in"^ (1.15 1.23 GPa).'3

The available data on the HT-S fibre composites remind us

that processing variables are important; in particular, those

variables that affect the void content. Detailed examination

of the available data showed that in 84 of the 89 reported

test series, the flexural strength of composites made with

HT-S fibre decHned markedly on cooling to 77 K. The five series for which strength increases were noted were reported

by Scheck to have occurred concomitantly with a change

to a vacuum-bagging method of fabrication. This suggests

that at least a pan of the observed drop in flexural strength on cooling might have been due to the presence of voids in the composites.

Continuing this with the other fibres, we observe on Fig.9a

that the order of decreasing strength is Modmor II. GY-70,

and Fibralloy 300. arri\ing finally at a group comprising the lowest reported strengths and consisting of Courtaulds HM,

HMG-25. and Thornel 50 fibres. The flexure strength of

this latter group is only about a third that of the HT-S

fibre composites. Among these other fibres, most data were

available for the high-modulus GY-70. for which the spread

of values at 77 K was 85-133 x lO^ lb in'^ (586-917 MPa);

the highest 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 Fig.9b suggests that there are

3O0

250

^00

OC 50 200

TemperotufB, K

Fig. 10 Ultimate flexural strength, a^, of miscellaneous advanced

composites

1 - boron-epoxy_(SP-272) ^''.2- HT-S/pol^mide 3 - HT-S/

epoxy-phenoloc ^

; 4 — GY-70/polyimide

xlO 50r 40-

-nxlO ^30

3C20

—— Medium modiius ' -2C

Low modulus

10

:5 >

OC 50 200

Temperoru'e, K

rf

23: 300

Fig.1 1 Initial flexural modulus, f j, of advanced composites:

uniaxial longitudinal

1 - Thornei 50; 2 - HMC-25; 3 - Modmor II (all reference 8)

CRYOGENICS . DECEMBER 1975 27 709
Table 2. Comparison of ranking of specific uniaxial graphite-epoxy composites by longitudinal flexure strength

and by longitudinal tensile strength (77 K)

Composite a^", 10^ lb in-2 Rank a'", 10^ lb in"2 Rank

A 178 1 147 1 B 172 2 87 6

C 167 o 1 no 1 Uo 2 D 146 4 45 8

c 117 5 92 4•t F 58 6 91 5 G 55 7 71 7 H 54 8 93 3

differences in the fibre-matrix bond strength between these two fibre types, the high modulus fibre having the poorer

bonding. Note that X-904 and 1004 epoxies were used with

both fibre types. Before leaving the data of Fig.9, it is of interest to consider whether or not the literature data support 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 Fig.9a with that of Fig. la 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 fibre and matrix)

for which data were available. The data are summarized

in Table 2. Results suggest that the flexure test results are not a very good measure of the relative tensile strengths of graphite/epoxy composites at 77 K. Insufficient data

existed for similar comparisons with other mechanical

properties. Rexure data on other types of advanced composites were somewhat meagre. Fig. 10 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, although still lower than that deve- loped 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 (see

Fig.2a). Data were available for two polyimide composites, one with the high modulus GY-70 fibre and one with the low modulus HT-S fibre. Of these, the HT-S composite de- veloped the higher strength. A small decrease in flexure strength is evidenced in these composites at 77 K, but the magnitude of the decrease is less than that reported for the uniaxial tensile strength. Finally, the HT-S/epoxyphenolic (HT424 Primer) appears to possess acceptable

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 (see Fig.2a).

Very little data were available on the flexural modulus.

Data for the longitudinal mode appearing on Fig. 1 1 reflect the modulus of the reinforcing fibres and appear to agree reasonably well with the moduli of the same composites in tension. The Thornel 50 data are an average of data from two matrix types, a modified ERL 2256 epoxy and NASA

Resin 2.^ A slightly higher modulus was reported for the

latter at both 295 K and 77 K.

Composite interlaminar shear strength Interlaminar shear strength is the resistance to failure be- tween layers in a layered laminate along the plane of their interface. Interlaminar shear is sometimes measured by the

guillotine method, in which shear is forced by the imposition of opposing but offset cuts across the width of a flat tensile specimen. Alternatively, the short-beam method

may be used. Here, the specimen resembles a flat flexural specimen but is proportioned to fail by shear on the central

layers of the composite. The guillotine method is reputed

to produce less scatter. However, the values obtained are usually higher than those obtained with the short-beam method. The latter method is the most widely used. A

modification of the short-beam test is used with filament- wound NOL ring specimens. Here, a short section of the

13 15 17 19

10 lb in

10 lb in

Fig.12 Histograms illustrating the reported range; of longitudinal interlaminar shear strength, a*

a - as reported for HT-S/X-904 graphite -epoxy at 77 K and

b — the reported changes in interlaminar shear strength in the same

composite upon cooling from 295 K to 77 K

710 28 CRYOGENICS. DECEMBER 1975 jf

.1
ring (concave downward) is substituted for the flat speci- ment. Results of the NOL segment and the flat short-beam

tests are not equivalent.

A wide range of interlaminar shear strength values may be

obtained with a single composite type. Fig. 12a is a frequency

histogram illustrating the range of longitudinal interlaminar

"shear strength reported in the literature for the graphite- reinforced composite HT-S/X-904 at 77 K. The substantial amount of data available for this composite reflects a comprehensive study of the effect of processing variables by

Maximovich and Scheck.'^ These data suggest that the maximum interlaminar shear strength obtainable with this type composite is on the order of 17-18 x 10^ lb in'^ (117-124 MPa).

The data of Maximovich and Scheck may be analysed fur- ther 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 Fig. 12b, we observe that the reported change in this para- meter 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. However, a more detailed look at the data shows

that those composites at the high end of the strength range

increased their interlaminar shear strength by about 4 000 lb in-2 (27 MPa) between 295 K and 77 K, suggesting

that the temperature dependence is affected by composite

quality.

Cryogenic interlaminar shear strength data have been published for other graphite-epoxy composites and for boronepoxy, PRD 49-epoxy, and for boron-aluminium.

However, available data are insufficient for statistical examination. One finds that the reported 77 K interlaminar shear

strengths for graphite-epoxy composites made with ten fibre types other than HT-S range from about 4-18 x 10^ lb in'^ (27-124 MPa). that is, covering about the same range covered by the HT-S data of Fig. 12a. Somewhat surprisingly, available data on boron-epoxy '^-^^ and boron-aluminium composites suggest that the maximum interlaminar shear strengths of these composites are almost the same as for the best of the graphite-epoxy composites, that is, about 18 x 10^ lb in^ (124 MPa) at 77 K

with about a 4000 lb in"^ (27 MPa) increase between 295 K

and 77 K. The only data on the interlaminar shear strength of PRD

49-epoxy composites are those of Hoggatt ^^•^'^ who report- ed NOL-ring segment values of approximately 400 lb in'^ (27 MPa) in both ERLB 4617 and NASA Resin 2 matrices. Values obtained with the latter of these were found to be

consistently higher by 40—70%. Very little temperature

dependence was observed. It is perhaps noteworthy that

the highest shear strengths reported for the graphite-epoxy

composites, 21x10^ lb in'^ (145 MPa), was also obtained

with NASA Resin 2 epoxy.^°

Composite bearing strength

Tlie only data found in the literature on the temperature

dependence of bearing properties of advanced composites

in the cryogenic range were published by Hertz et al for boron and Borsic-reinforced 6061 aluminium, including the

steel and titanium-reinforced hybrids. These data, appearing

200 250 300

Temperoture, K

Fig. 13 Ultimate bearing strength, o^", of boron-aluminium

composites, uniaxial longitudinal direction

1 — 4.2 mil Borsic-aluminium (6061 ); 2 — 5.6 mil Boron-aluminium

(6061); 3 - 4.2 mil Borsic-steel-aluminium (6061); 4 - 4.2 mil Borsic-titanium-aluminium (6061) (all reference 13)

on Fig. 13, are for bearing strength in the unaxial longitudinal direction, that is, the stress required for bearing pull- out, rather than the more conventional bearing yield

strength. The data on Fig. 13 refer to m.aterial hole-to-free surface dimensions two and four times the diameter of a No 10 steel pin (~ 4.8 mm). An increase in bearing strength

with decreasing temperature is observed for both conventional boron and Borsic-reinforced materials, the effect being more pronounced in the 4D tests. The Borsic-steel- aluminium hybrid performed much like the conventional

materials, while the Borsic-fitanium-aluminium hybrid per- formed somewhat erratically, although the highest strengths were developed with this latter composite.

Dynamic mechanical properties

Composite fatigue

Tlie only available data on the temperature dependence of

the fatigue properties of advanced composites were gener- ated during cyclic pressure testing of filament-wound

pressure vessels. As the composites in these vessels were

cross phed and subjected to biaxial stressing, the data are of value only in a comparative sense. Hansom has studied the cyclic fatigue perfermance of

pressure vessels overwrapped with Thornel 50 graphite

fibre using an ERL 2256/ZZL 0820 epoxy 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- fibre reinforced vessels. Hanson concluded that, on the

basis of specific strength, graphite-fibre reinforced vessels would have fatigue performance superior to that of glass- fibre reinforced vessels after only 80 cycles.

CRYOGENICS. DECEMBER 1975 29 711
Alfring et al have reported cryogenic cyclic fatigue data on pressure vessels filament-wound with 4 mil boron in a Polaris-resin matrix. 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.

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 falhng to 40-50% of single-cycle values after 10 000 cycles. This is about the same as for

glass reinforcement. These results suggest that graphite may be superior to boron

as a reinforcement fibre for composite structures subject to

fatigue at cryogenic temperatures. However, these data must be considered very tentative in view of the afore- mentioned tendency for the 4 mil boron fibre to split longitudinally when subjected to transverse stresses such

as are present in biaxially-loaded pressure vessels. The

fatigue performance at cryogenic temperatures of composites reinforced with 5.6 mil boron fibre should be substantially better than with the 4 mil fibre, if fibre splitting is a factor in the smaller diameter fibre.

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^ boron-epoxy material tested transverse to the uniaxial longitudinal direction. These data indicate a slight increase in impact strength from 22.8 ft lb in'' (121.7

N m m-i ) at 295 K to 27.0 ft lb in' (144 N m m"' ) at 20 K

Concomitantly, a slight decrease from 46.3 to 39.5 ft lb in'' (247 to 211 N m m'' ) was reported for unnotched speci- mens over this temperature range. The notched/unnotched

ratios of about 0.49 at 295 K and about 0.68 at 20 K in- dicate some notch sensitivity. Sumner and Davis ^ have

reported a 12 ft lb in'' (64 Nmm"') Charpy V-notch

strength at 295 K for 25 v/o stainless steel wire reinforced 2024 aluminium, this value increasing to 18 ft lb in'' (96 N m m'' ) at 77 K. Cryogenic impact strength data were not available for the other types of advanced composites. However, room temperature data indicate that the graphite-reinforced composites have substantially lower impact strengths than boron-reinforced types. 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 1 for glass- reinforcement, where notched values transverse to the

uniaxial longitudinal direction were seen to range from

67-162 ft lb in' (357-865 Nmm"') at 77 K.* The

critical factor has been identified as the stress-strain be- haviour of the fibre reinforcement, higher impact values being associated with higher failure strains of the fibre whenever tests are made in the longitudinal (cross-fibre fracture) mode.'^^''^'^ Impact strengths of uniaxial composites tested in the trans- verse direction (anvil impact parallel to the fibres) 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 possibihty of an overall improvement of composite toughness. Larsen ^ has reported efforts to improve

Temperature, K

Fig. 14 Longitudinal thermal expansion, £L/L, of uniaxial advanced composites

1 -Thornel 75/ERLB4617 ^

' ; 2 - G Y-70/X-904

3 - Fibrailoy 300/X-904 '^•4 - HT-S/X-904 5 - PRD49-1/

ERLA4617 ".^°;6 - PRD49-I/NASA Resin 2 ; 7 - boron/ epoxy SP-272; 8 — 5.6 mil boron/alumininium 6061 ; 9 — 4.2 mil Borsic-titanium-aluminium 6061 ; 10 — 4.2 mil Borsic-steel- aluminium (6061); 11 —4.2 mil Borsic/aluminium (6061) (all reference 13)

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, that is, 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 per- formance at cryogenic temperatures. Larsen concluded

that improved overall cryogenic properties would be ob- tained by use of the NASA Resin 2 epoxy formulation with

graphite-reinforced composites.

It appears at this time that substantial improvement of impact strength of the advanced composites will require deve- lopment of hybrids, possibly combining glass with the advanced fibres.

Thermal properties

Composite thermal expansion and contraction Unlike the glass-reinforced composites, some advancedfibre composites expand in the fibre direction when cooled

to cryogenic temperatures. Available data for the temperature dependence of dimensional changes in the uniaxial longitudinal direction are summarized in Fig. 14, while comparable data for the uniaxial transverse direction appear

in Fig. 15.

712

30

CRYOGENICS.. DECEMBER 1975
The graphite-reinforced composites display a ver> small longitudinal expansion on cooling, slight!)- larger expansions

being reported for the high modulus Thomel 75 and GY-70

fibre composites than for composites made with the mediummodulus Fibralloy 300 and the low-modulus HT-S fibres. The PRD 49 (Kevlar 49) composites undergo a comparative!)

large longitudinal expansion on cooling, reportedly reaching

a ma.ximum of about 8—10 x 10"^ at about "~ K.^^'^°

Slightly larger expansion was reported with a NASA Resin 2 matrix than with ERLA 46 1 7. The other advanced composites undergo linear contraction on cooling. The least longitudinal contraction was reported for the commercial

SP-272 boron-epo.xy product, while boron-aluminium undergoes the largest contraction. The matrix contraction properties appear to be dominating in the latter composites, as re- lativeK- little difference is seen among the four variants, in- cluding the hybrids. The 5.6 mil composite displays a sUghtly lower contraction than does the 4 mil product. The longi- tudinal thermal contractions of these boron-aluminium

composites are about 25?c of that of an unreinforced 6061

aluminium 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 obsened in the fibre direction,

as witnessed by the necessity to change the ordinate scale on

Fig.15. 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 graphite-reinforced composites appear to underxi0"'«

<

-

-160 r

0 5C 100 ;50 2GC ZzZ iZC

Temoeroture, K

Fig.15 Transverse thermal expansion, AL/L, of uniaxial advanced

composites

1 -boron/epoxy (SP-272) 2 - Fibralloy 300 ERLA 4617

HT-S/X-904 '^;3 -4.2 mil Borsic-titanium/aluminium 6061 ; 4 - GY-70/X-904 '^;5 - boron/aluminium 6061 '^;6 -4.2 mil Borsic-steel/aluminium (6061) 7 - PRD49-I/ERLA 4617

0 2'C 50 iOO ioO 2X 25G 30C

j

Fig. 16 Thermal conductivities, \, of advanced composites

a - Uniaxial longitudial; 1 - HTS/X-904 2 - Thornel 50-Polaris 3 - B-epoxy, SP-272 ; 4 - B-epoxy, 4.0 Boron-Polaris j 5 - PRD 49-epoxy ^ | b - Unixiai transverse:J - HT-5 X-904 '-\-2 - Thornel 50 - Polaris^'; i 3 - B-epoxy, SP-272 j

I go about twice the transverse thermal contraction of the | boron-reinforced composites, reflecting the lesser constraint provided in this direction by the small graphite fibres. .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 fibre volume fraction. The conventional boron 6061 aluminium 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 the boron-aluminium | composite is almost completely matrix dominated. No sign- | nificant difference in thermal contraction were reported for 4.2 mil Borsic and 5.6 mil boron fibre composites.

Hie PRD 49-epoxy composites undergo the largest transverse thermal contraction. Comparing the temperature dependence

of PRD 49-1 NASA Resin 2 (curve 8) with that reported by

Soffer and .Molho - for NAS.A Resin 2 alone (dashed curve), indicates that the matrix is also dominating the transverse dimensional changes in tliis composite. The PRD 49 fibres do not offer appreciable restraint to transverse contraction of the matrix because, as seen in Fig. 14, this fibre expands | significantly in the longitudinal direction during cooling. Of

necessity, this is accompanied by signitlcant transverse fibre contraction. Cooling to cr\-ogenic temperatures would be

expected to produce sizable interfacial shear forces betw^een

the PRD 49 fibre and matrix, with lesser forces normal to

the fibres.

Composite thermal conductivity

Few- data were available on the temperature dependence of

uniaxial longitudinal and uiua.xial transverse thermal conducti\it\- of advanced composites. Those which were avail- able are plotted in Fig. 16. WTien considering these data, the

j

I CRYOGENICS . DECEMBER 1975 31
Temperature , K

Fig. 17 Specific heat, Cp, of advanced composites

1 - HT-S/X 904, GY-70/X 904 '^• 2 - B-epoxy (SP-292) 3 — boron/6061 aluminium ; 4 — Borsic-titanium/aluminium

5 — Borsic-steel/aluminium ; 6 — graphite/phenolic

reader should be aware of the difficulties and potential sources of error in determining thermal conductivity within

the cryogenic range. The comments made in Part 1 of this paper on this subject apply equally well here. Furthermore,

the volume fraction of fibre will infiuence the conductivities,

particularly in graphite-reinforced composites. For these

reasons, the data on Fig. 16 should be considered as only

indicating trends. The graphite-reinforced composites have relatively high thermal conductivities in the fibre direction, refiecting re- latively high conductivity of the graphite fibres compared

to epoxy. Longitudinal conductivity in these materials is highly temperature dependent. The boron-reinforced composites have much lower conductivities than do the graphite- reinforced composites in the fibre direction and show a comparatively small temperature dependence of thermal conductivity. Data were not available for boron-reinforced aluminium; however, the thermal conductivity of such materials wall certainly be much higher than that of a polymericmatrix composite. The conductivity should be approximately, half that of unreinforced aluminium in a typical 50

volume % boron-aluminium composite. Transverse thermal

conductivities are matrix dominated and very low for both

the graphite and boron-reinforced epoxy materials. The only available data for PRD 49 indicate a longitudinal thermal conductivity approximately the same as that for boron-epoxy within the cryogenic range. These data suffice to show that the thermal conductivities of

ihe epoxy-matrix advanced composites, particularly boronreinforced, are low enough to make them attractive for cryogenic structural components.

Composite specific heat

Fig. 17 shows the specific heat of the advanced composites

to be similar to those of the glass-reinforced composites dis- cussed in Part 1 of this paper. Again, an almost linear temperature dependence of Cp is observed from 295 K to 77 K. As a group, the specific heats of the composites are

slightly lower than that of aluminium for much of the region

between 295 K and 77 K, but substantially higher than that

for copper and most other metals. The temperature depen- dence of Cp for aluminium, copper and titanium are included on Fig. 17 for comparative purposes. The highest specific heat at cryogenic temperatures was re- ported for boron-aluminium, the aluminium matrix undoubtedly contributing significantly to this relatively high value. The depicted curve is the average of data reported by Hertz

et ap3 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-aluminium hybrids had slightly lower specific heats, probably reflecting the contributions from the steel and titanium reinforcements. The specific heat of the boron-epoxy composite lies between that of the boron-aluminium 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 Fig. 17, 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 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 Cp 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 under- standing of the present state of knowledge as to the magni- tude and temperature dependence of the properties of ad- vanced fibre composites at cryogenic temperatures. It is apparent that present knowledge is incomplete in many re- spects; nevertheless, the available data justify considerable optimism as to the future applicability of these materials

as structural components operating at cryogenic

temperatures.

Available data indicate that commercial state-of-the-art boron-reinforced epoxy and aluminium perform well at cryogenic temperatures. The higli tensile and compressive

strengths and elastic moduli of these materials remain constant or improve as the temperature is lowered. Equally

impressive is the relatively small scatter in the data, suggest- ing 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. Although actual data are not available, boron-aluminium

composites are certain to have a much higher thermal conductivity than any polymer-matrix composite (approximately

half that of the unreinforced alloy). This is a distinct dis- advantage in many cryogenic applications. On the other hand, the transverse strength and modulus of boron- aluminium is much higher than that of polymer-matrix

composites. The latter is a distinct advantage of metal- matrix composites, simplifying composite design and in- creasing component reliabihty.

Available data on the cryogenic performance of the graphite- reinforced epoxy class of composites are less convincing.

714 32 CRYOGENICS. DECEMBER 1975
! These materials can be produced with moduh higher than

' those obtainable with boron-reinforced materials. Graphiteepoxies have the highest thermal stability of any composite

I class, which is an advantage for many cryogenic applications,

j However, the ultimate tensile strengths of the graphiteI epoxies are substantially lower than for composites with

boron reinforcement, and the strength tends to decrease

' ~upon cooling to cryogenic temperatures. Mechanical proI perty test resuhs are frequently unpredictable and contraj dictory. It is not at present known how much of the erratic behaviour is due to problems in testing and how much is an

inlierent characteristic of the material. Certainly, the overI all performance of graphite-fibre overwrapped pressure

I vessels reported by Larsen and Simon '° was much better

i than would have been predicted from their basic test data. We therefore conclude that graphite-reinforced composites

i have potential for structural use at cryogenic temperatures,

in particular, for applications requiring dimensional

stability. However, further development and testing are required.

PRD 49 (Kevlar 49)-reinforced epoxies also appear pro- mising for cryogenic applications, due to the relatively low

I cost of the fibre and the improved modulus compared to

glass. However, present data on this material are minimal, and further investigation must be undertaken before this

i material can be appUed with confidence in a cryogenic en- vironment. PRD 49 is unusual in that it undergoes a signj ficant expansion in the fibre direction during cooling to cryogenic temperatures. Large residual stresses between the

I fibre and the epoxy matrix might be expected to adversely

j affect the mechanical properties of composites made with

; this material. Nevertheless, the work of Hoggatt ^^'^^has shown that, with proper design, PRD 49-epoxy composites

' can be used to good advantage in some cr>'0genic applications. As no compressive strength data were available at cryogenic temperatures, the potential user should investigate

this parameter in applications where compressive loading is a factor.

In contrast to the boron-reinforced materials, the available data suggest that both graphite and PRD 49-reinforced

II composites may benefit from matrix optimization when

' used at cryogenic temperatures. Larsen and Simon have recommended NASA Resin 2 for both uniaxial and crossply

graphite-reinforced composites. The work of Hoggatt

suggests that this type of matrix is also satisfactory for PRD

.i 49 reinforcement. ^^'^^ When considering the use of NASA

I Resin 2, the reader should be aware that this resin has been

optimized for cryogenic use by the addition of flexibilizers, which significantly reduce the strength at elevated temper-

! atures. In particular, care must be taken to properly sup1 port components made with this resin whenever elevated

If temperature baking of a component is required.

The foregoing is not intended to imply that the NASA Resin

2 formulation is clearly the best resin for all cryogenic

I applications. Composite manufacturers may well recomji mend other types of epoxies for cryogenic applications. The

: present report is not intended to be prejudicial to such recommendations. However, the purchaser is advised to

|| request documentation that the recommended resin has

[j shown itself to be suitable for the intended cryogenic applicI ation. In particular, the suitability of various resins for

!: crossply composites subjected to fatigue loading under

cryogenic conditions has not been adequately investigated.

Summary of primary property trends in uniaxial ' advanced composites upon cooling to cryogenic | temperatures

j

Tensile strength i Boron-epoxy and boron-aluminium show little temperature

dependence of strength. The same appears to be true for ; PRD 49 (Kevlar 49)-epoxy, although data are minimal.

The graphite-epoxies tend toward decreased strength on

cooling, 77 K values being of the order of 20 x 1 0^ lb in'^ { (138 MPa) lower than at room temperature. "j

1

N Tensile modulus | Boron-epoxy, boron-aluminium, and graphite-epoxy com- \ posites appear to retain their room-temperature moduli

j when cooled to 77 K. Minimal available data indicate that J PRD 49 (Kevlar 49)-epoxy composites undergo a substan- l tial modulus increase on cooling below 77 K.

|

Compressive strength \ Compressive strengths of the advanced composites tend to ^ either increase (boron-epoxy, graphite-epoxy) or to remain

essentially unchanged (boron-aluminium) during cooling '

to 77 K. Compressive strength data were not available for -i PRD 49 (Kevlar 49)-epoxy composites. i

Compressive modulus i Compressive moduli generally remain unchanged upon coolj ing to cryogenic temperatures, except for a small moduli I ijicrease reported for boron-epoxy. Data were not available i for PRD 49 (Kevlar 49)-epoxy.

j Ultimate tensile strain ] Available data indicate a small (5—15%) decrease in longi- i tudinal strain at faUure on cooling boron-aluminium and

graphite-epoxy to 77 K; the latter being the more temperature sensitive. Transverse failure strain appears almost independent of temperature in these composites. Data were

not available for PRD 49 (Kevlar 49) or boron-epoxy.

In-plane shear strength and modulus

Insufficient data are available to indicate trends in these

properties. One reference cited about a 14% drop in shear

strength, concomitant with about a 50% shear modulus ^ increase, for a graphite-epoxy composite upon cooling to 77 K; however, variations in fibre type and matrix com- i position will undoubtedly affect the temperature sensitivity j of these properties. No data were available for other advanced composites.

\ Thermal expansion and contraction Graphite-epoxy is the most dimensionally stable of the advanced composites, undergoing a very sUght expansion

longitudinally (fibre direction) and moderate transverse contraction on cooling. Boron-epoxy is the next most

stable, showing moderate contraction in both directions on cooling. Boron-aluminium contracts about twice as much as boron-epoxy longitudinally; however, the contraction is only about 20% of that of the unreinforced j metal. Transverse contraction of boron-aluminium is al- '

CRYOGENICS . DECEMBER 1975 33 715.
most the same as that of the unreinfcreed metal. PRD 49

(Kevlar 49)-epoxy composites are the least dimensionally

stable, expanding substantially in the longitudinal direction and contracting substantially in the transverse direction.

Thermal conductivity

Boron-epoxy and PRD 49 (Kevlar 49)-epoxy have substantially lower longitudinal thermal conductivity than graphiteepoxy composites. Differences in conductivity decrease as temperature decreases. Conductivity in the transverse direc- tion is approximately the same for graphite and boron-epoxy,

being much lower in this direction. Data were not available

for PRD 49 (Kevlar 49)-epoxy and for boron-aluminium.

Specific heat The specific heat of the advanced composites appear to be

almost linear from 295 K to 77 K. As a group, the values

are slightly lower than that of aluminium, but substantially

above that of titanium or copper. The specific heat of boron-aluminium is slightly higher than that of boron- epoxy or graphite-epoxy. No data were available for PRD

49 (Kevlar 49)-epoxy.

The author thanks Dr R. P. Reed for consultation and review of the completed manuscript. Many thanks are also due

to 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.

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 re- commendation by NBS. Generic names have been substit- uted whenever it was possible to do so without sacrificing

clarity. Manufacturers of the trade named materials are

listed in the Appendix.

Appendix

The following materials are referred to in this report:

Fibres Graphite

HT-S, HM-S

HMG-25, HMG-50

Modmor I, II Thornel 25, 50

Samco 320

Fibralloy 300

GY-70

Boron

Borsic

Others PRD 49 (Kevlar 49)

Resins Epoxies

X-904, X-915

ERL2256, ERLB 4617

Epon 828, 1031, 58-68R

Courtaulds Ltd

Hitco Corp

Morganite Ltd Union Carbide Corp

Samco Corp Monsanto Corp

Celanese Corp

Hamilton Standard Corp

E. I. DuPont de Nemours,

Inc

Fiberite Corp

Union Carbide Plastics Co

Shell Chemical Corp,

Plastics & Resin Div

US Polymeric Corp

Hercules Corp

Monsanto Corp

American Cyanimid

Union Carbide Plastics Co

Emery Industries, Inc

Minnesota Mining &

Manufacturing Co

AVCO Corporation

E-350, E-787

Ciba 8183/137, 3002

Others

Skybond 703 (polyimide)

HT424 (epoxy-phenolic)

Flexibilizers, Hardeners

ZZL 0820

Empol 1040

DSA - dodecenyl succinic anhydride

BDMA — benzyldimethylamine

NMA — nadic methyl anhydride

Boron Composite Products SP-272

5505

Bibliography

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Facility: California

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(1962) (AD 288 944)

1.1 Brink, N. O. 'Mechanical behaviour of reinforced plastics at cryogenic temperatures', Technical Papers, 20th Annual

Technical Conference, Society of Plastics Engineers, Vol 10, Section XV-2 (1964)' 1-19

1.2 Brink, N. O. 'Mechanical behaviour of reinforced plastics

at cryogenic temperatures, Societv of Plastics Engineers Journal 20 (1964) 1123

1.3 Brink, N. O. 'Mechanical behaviour of reinforced plastics at cryogenic temperatures', Narmco Research and Development Report, Code No 105-4 (1964)

2 Chamberiain, D. W., Lloyd, B. R., Tennant, R. L. Deter- mination of the performance of plastic laminates at cryogenic

temperatures', ASD-TDR-62-794, Part 11 (1964) (N64-24212)

2.1 Chamberlain, D. W. 'Tensile fatigue testing at temperatures down to 20 K', Advances in Cryogenic Engineering 9 (1964) 131

2.2 Chamberiain, D. W. 'Mechanical properties testing of

plastic laminate materials down to 20 K\ Advances in Cryogenic Engineering 10 (1965) 117

Contract NAS 8-11070

Contractor: National Aeronautics and Space Administration, George C. Marshall Space Flight Center, Huntsville, Alabama

Research: Goodyear Aerospace Corporation, Akron,

Facility: Ohio

3 Toth, L. W., Boiler, T. J., Butcher, I. R,, Kariotis, A. H., Yoder, F. D. 'Programme for the evaluation of structural reinforced plastic materials at cryogenic temperatures', NASACR-80061 (Final) (1966) (N67-12051)

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) 1

716 34 CRYOGENICS . DECEMBER 1975
3.2 Toth, L. W., Kaiiotis, A,H. 'An assessment of test specimens and test techniques useful to the evaluation of structural reinforced plastic materials at cryogenic temperatures', Adi-nnces in Cry ogenic Engineering 10 (1965) 126

3.3 Toth. L. "Properties of glass-reinforced epoxy through the 20 K range". Modem Plastics 42 (1965) 123

3.4 Toth, L, W,, Burkley, IL 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 (1967)

3.5 Toth, L. W., BoUer, T. J„ Kariotis, A. H., Yoder, F. D. 'Programme for the evaluation of structural reinforced

plastic materials at cryogenic temperatures", NASA CR-64005

(1963-64) (N65-29724)

Contract NAS 3-6297

Contractor: National Aeronautics and Space Administration, Lewis Research Center, Qeveland, Ohio

Research .Aerojet General Corporation, Azuza,

Facility: California 4 Lewis, A., Bush, G. E. "Improved cryogenic resin-glass filament mound composites'. NASA CR-72163 (Final) (1967) (N67-31856)

4.1 Lewis, A., Bush, G. E., Creedon, J. "Improved cryogenic

resin glass filameni-wound composites", NAS.A Interim Report CR-54867 (1966) (N66-28040)

Contract NAS 6-6287

Contractor: National Aeronautics and Space .Administration, Research Center, Geveland, Ohio

Research Aerojet General Corporation, Azuza,

Facility: California

5 Soffer, L. M., Motho, R., "Cryogenic resins for glass filament- wound composites", NASA CR-72114 (Final) (1967) (N67-25076)

5.1 Soffa, L. M., Molho, R,, 'Mechanical properties of epo.xy resins and glass epoxy composites at crj'ogenic temperatures', Cryogeni

Properties of Polymers (J. L. Koenig (ed)] (Marcel Dekker, New

York, 1968) 87 (Identical to NASA CR-84451. 1967, (N66-27217)

Contract FO4701-69-C-0059

Contractor: Space and Missiles Systems Organization, Air Force Systems Command, Los Angeles Air Force Station, Los Angeles, California

Research The Aerospace Corporation, El Segundo,

Facility: California

6 Pepper, R. T., Rossi, R. E., Upp, U. W., Riley, W. E. 'Develop- ment of an aluminium-graphite composite'. SAMSO-TR-70-301

(1970) (AD 718 409)

6.1 Pepper. R. T.- Upp, J. W.. Rossi, R. C, Kendall, E. G. 'Aluminium graphite composites' SAMSOTR-70-174 (1970)

(.AD 706 883). (Identical lo Metal lurgical Trans 2 (1971) 117

6.2 Rossi, R. C. Pepper, R, T., Upp, J. W., Riley, W. C. "Develop- ment of aluminium-graphite composites' Ceramic Bull 50

(1971) 484

Contract NAS 8-11508

Contractor: National Aeronautics and Space Administration, George

C. Marshall Space Flight Center. HuntsvUle, Alabama

Research Harvey Engineering Laboratories, Torrance,

Facility: California

7 Summer, E. V., Davis, L. W. 'Development of ultrahigh strength, low density aluminium sheet and plate composites', NASA

CR-85863 (Final) (1966) (N6731181)

7.1 Davis, L. W. 'Composites at low temperature', Paper No. 15, Seventieth Annual Meeting of the American Society for Testing

Materials (1967)

Contract NASA DPR C 10360-B

Contractor: National Aeronautics and Space Administration, Lewis Research Center, Cleveland, Ohio

Research Naval Ordnance Laboratory, Silver Springs,

Facility: Maryland

8 Simon, R, A., Alfring, R. 'Properties of graphite fiber composites at cryogenic temperatures', NASA CR-72652 (NOLTR 69-183) (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) (1971) Task III (AD 882 972)

10 Larsen, J. V., Simon, R, A. 'Carbon fiber composites for cr>'ogenic filament-wound vessels', NASA CR-1 20899 (NOLTR 71-201) (1972) Tasks IV, V, and VI (N73-11553)

10.2 Larsen, J. W. 'Fracture energy of CBIN/epoxy-carbon fiber composites', Technical Papers, 26th Annual Technical, Conference. Society of the Plastics Industry, Section 10-D

(1971)

Contract NAS 8-26198

Contractor: George C. Marshall Space FMght Center, HuntsviUe, Alabama

Research General Dynamics Convair, San Diego,

Facility: California

1 1 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 (1974)

12 Scheck, W. G. 'Development of design data for graphite

reinforced epoxy and polyimide com.posites', 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., Scheck, W. G. 'Development of graphite/ polyimide composites', National-Technical Conference,

Society of Aerospace Material and Process Engineers, Vol 3 (1971) 717

Contract F33615- 70-1442

Contractor: Air Force Materials Laboratory, Wright-Patterson

Air Force Base, Ohio

Research General Dynamics/Convair, San Diego,

Facility: California 13 Hertz, J., Christian, J. L., Varlas, M. 'Advanced composite

applications for spacecraft and missiles. Phase I Final Report, Volume II: Material Development', AFML-TR 71-186,

Vol 2 (1972) (AD 893 715L)

13.1 Forest. J. D.. Fujimoto, A. F., Foelsch, G. F. 'Advanced

composite applications for spacecraft and missiles. Phase I Final Report. Volume 1 : Structural Development". .AFMLTR-71-186. Vol 1, (1972)

13.2 Forest. J. D., Varlas, M. 'Advanced composite applications

for spacecraft and missiles. Final Report, AFML-TR-72-278

(1973)

13.3 (Thristian, J. L., Campbell, M. D. 'Mechanical and physical

properties of several advanced metal-matrix composite

maxeiials'. Advances in Cryogenic Engineering 18 (1973) 175

2. Thermophysical properties

Contract AF-33 (657)-9160

Contractor: Air Force Materials Laboratory, Wright-Patterson

Air Force Base, Ohio

CRYOGENICS . DECEMBER 1975 35 717
Research General Dynamics/Astronautics, San Diego,

Facility: California

The following reports are in a series entitled 'Therniophysical

properties of plastic materials and composites to liquid hydrogen

temperature (-423°F)'

14 Haskins, J. F.; Campbell, M. C, Hertz, J., Percy, J. L. ML-TDR-64-33. Part 1 (1964) (AD 601 337)

1 5 Campbell, M. D., Hertz, J.. O'Barr, H. L., Haskins, J. F. ML-TDR-64-33, Part 2 (1965) (X65-18921)

16 Campbell, M. D., O'Barr, G. L., Haskins, J. F., Hertz, J. ML-TDR-64-33, Part 3 (1965) (AD 468 155)

16.1 Hertz, J., Haskins, J. F. 'Thermal conductivity of reinfroced

plastics at cryogenic tempeiatmes'. Advances in Cryogenic

Engineering 10 (1965) 163

16.2 Campbell, M. D. Thermal expansion characteristics of some plastic materials and composites from room temperature

to -253°C' Advances in Cryogenic Engineering 10 (1965) 154

16.3 Campbell, M. D., Haskins, J. F., O'Barr, G. L., Hertz, J. 'Thermophysical properties of reinforced plastics at cryogenic temperatures' Journal of Spacecraft 3 (1966) 596

(See also reference 13)

Contract F33615- 73~C1388 (Work currently in progress)

Contractor: Air Force Materials Laboratory, Wright-Patterson

Air Force Base, Ohio

Research General Dynamics/Convair, San Diego,

Facility: California

1 7 Forest, J. D., Schaeffer, W. H. 'Advanced composite missile and space design data'. General Dynamics Report GDCACHB72-001-1, Progress Report No 1 (1972)

18 Forest J. D. 'Advanced composite missile and space design

data', General Dynamics Report GDCA-CHB72-001-2,

Progress Report No 2 (1973)

3. General bibliography

19 Aleck, B. 'Fibreglass-overwrapped 2219-T87 aluminium

alloy low-pressure cryogenic tankage'. Society of Aerospace

Material and Process Engineers National Technical Conference, Space Shuttle Materials, Vol 3 ( 1 97 1 ) 1 3

1

20 AJfring, R. J., Morris, E. E., Landes, R- E. 'Cycle-testing of boron fUament-wound tanks', NASA CR-72899, National Aeronautics and Space Administration, Lewis Research

Center A ( 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 (1969) (N70-25098)

22 Baucom, R. M. 'Tensile behaviour of boron filament- reinforced epoxy rings and belts', NASA TN D-5053,

Langley Research Center, Hampton, Virginia ( 1 969) (N69-19918)

23 Benton, W., Carr, R., Cohen, A., Gustafson, G., Lankton, C.

Zeldin, B. 'Propellant storability in space', RPL-TDR-64-75

(Final), Air Force Systems Command, Edwards Air Force

Base, CaUfornia, (1964) (AD 603 215) 24 Brechna, H. 'Superconducting magnets for high energy physics

applications', Proc ICECl (Heywood Temple Industrial Publishers Ltd, London, 1968) 119 (CFSCl N67-36009)

25 BuUard, B. R. 'Cryogenic tank support evaluation', NASA

CR-72546, NASA Lewis Research Center, Cleveland, Ohio

(1969) (N70-13085)

26 Campbell, M. D. 'Development of thermal expansion capabilities and the investigation of expansion characteristics of

space vehicle materials'. General Dynamics/Astronautics

Report ERR-AN-251 (1962)

27 Campbell, M. D. 'Development of the thermal expansion

capabilities and the investigation of the thermal expansion

characteristics of space vehicle materials (11)', General Dynamics/Astronautics Report ERR-AN-450 (1963)

28 Caren, R. P., Coston, R. M., Holmes, A. M. C, Dubus, F. 'Low-temperature tensile, thermal contraction, and gaseous

hydrogen permeability data on hydrogen-vapor barrier

materials'. Advances in Cryogenic Engineering 10 (1965) 1 71 29 Chiao, T. T., Moore, R. L." 'Tensile properties of PRD-49

fiber in epoxy matrix', J Comp Mat 6 (1972) 547

30 Cooper, G. A., Sillwood, J. M. 'Multiple fracture in a steel reinforced epoxy resin composite', 7 Maf Sci 1 (1972) 325

31 Daiwish, F., Tetelman, A. S. 'Mechanical behaviour 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., Zender, G. W. 'Mechanical behaviour 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, 1 (1962) 34 Fontana, M. G., Bishop, S. M., 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, (1953) (AD 27726)

, 35 Freeman, S. M. 'Properties of vapour barriers, adhesives and foams at cryogenic and elevated temperatures', Lockheed

Aircraft Corporation Report ER-5687 (1962) 36 Freeman, W. T., 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) 121 37 Funk, C. W., Dixon, C. E. 'Cryogenic radiation damage in structural polymers', Trans of the American Nuclear Society 9 (1966)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 (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 (1971) (N71-22401)

40 Gray, P. D., Cornelius, C K., O'DonneU, J. D., Howard, W. W.

'Rockets in space environment, Volume 1 : The experimental programme', RlD-TDR-63-1050, Aerojet General Corporation

(1963) (N63-20999)

41 Greer, F. 'Flexural properties of Conolon 506 at room

temperature ~320°F and -423°F', Convair/Astronautics Report 55E 522 J (1971) (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

(1969) (N69-35811)

43 Hall, J. 'Cryogenic tensile tests - epoxy fiberglass', Douglas

Aircraft Company Report MP 1348 (1961) 44 Hanson, M. P. 'Effects of temperature and creep characteristics of PRD^9 fiber epoxy composites', NASA TN D-7 1 20, National Aeronautics and Space Administration, Lewis

Research Center, Cleveland, Ohio (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 (1969) (N69-31300) (identical

to SAMPE 15, 249) 46 Hanson, M. P., Richards, H. T., 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 (1965) 47 Hanson, M. P. 'Glass-, boron-, and graphite-filament-wound

resin composites and liners for cryogenic pressure vessels', NASA TN D-4412, National Aeronautics and Space

Administration, Lewis Research Center, Cleveland, Ohio

(1968) (identical to NASA TM X-52350, 1967)

718 36 CRYOGENICS. DECEMBER 1975
48 Hanson, M, P. 'Static and dynamic fatigue behaviour of glass filament-wound pressure vessels at ambient and crj ogenic

temperatures'. NASA TN I>-5807. National Aeronautics and

Space Administration. Lewis Research Center, Cleveland. Ohio 1 1970) iCFSTl-CSCL-20 K)

49 Haskins. J. F., Hertz. J. Thermal conductivity testing of

coast F-224-6 phenolic-fiberglass laminate'. General Dynamics/

Convair Report No AR-592-:-482 (1963) 50 Haskins. J. F., Huilich. A. 'Measured values for the coefficients of linear expansion of Plycel 420 and Conolon 506

at low temperatures", Convair/ Astronautics Report .\1RG-154 (1960)

51 Haylett J. W., Rottmayer, E.. Butcher, I. 'Advanced

composite material study for millimeter wavelength antermas'. Technical Report AFML-TR-71-205, Vol 1, Aii Force

Materials Laboratory. AVriahl-Patterson Air Force Base, Ohio

(1971) (.AD 893 368)

52 Haylett C E, 'Advanced composite material study for millimeter wavelength antennas. Volume II. Environmental

Tests'. AFML-TR-7'l-205-Vol 2, Air Force Materials Laboratory, Wright-Patterson .Air Force Base. Ohio (1971) (AD 893 3^58 L) 53 Herring, H. W.. Baucom. R, M.. Pride. R. A. "Research on

boron filaments and boron reinforced composites', 10th

National Symposium, Society of Aerospace Material and

Process Engineers. Vol 10 (1966) B-21 to B-34

54 Hertz, J. 'TensUe testing of Conolon 506 at room and sub- zero temperatures', Convair/ Astronautics Report MR(3-120

55 Hertz. J. "Tensile testing of .Adlock 85 1 . .Adlock PG-L.A and

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

Astronautics Report MRG 237 (1961)

56 Hertz, J. 'Investigation of potential low temperature insulations' General Dynamics/ Astronautics Report GS/.A-ERR-.A.N-668 (1964)

57 Hertz, J. 'The effect of cryogenic temperatures on the mechanical properties of reinforced plastic laminates'. General Dynamics Report AR-592-M15 (1963) (AD 405 170)

58 Hertz, J. "Investigation into the high-temperature strength degradation of fiber-reinforced resin composite during ambient aging". General Dynamics/Convair Report No

GDCA-DB'g73-005 (Final'Contract N.AS 8-27435) (1973)

59 Hoggati, J. T. 'Development of cryogenic PRD-49-1 filament- wound tanks". NASA CR-120835, National Aeronautics and

Space Administration, Lewis Research Center, Cleveland, Ohio (1971) (N72-24941)

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

for pressure vessel applications'. Society of Aerospace

.Material and Process Engineers National Technical Conference, Space Shuttle .Materiais."^Vol 3 (1971) 157

61 Hust J. G. "Low temperature thermal conductivity of two

fibre-epoxy composites' Cryogenics 15 (1975) 126 62 Johnston. H. L.. Brooks. H. E. 'Impact strength of various metals at temperatures down to 20° absolute'. Ohio State

University Cryogenic Laboratory Report TR 264-17 (1952)

63 Keller. C W. "Fiberglass supports for cryogenic tanks', N.ASA CR-1 20937 (Final), National Aeronautics and Space

Administration. Lewis Research Center, Qeveland. Ohio (1972MN72-33564)

64 Keriin, E. E., Smith, E. T. ".Measured effects of the various combinations of nuclear radiation, vacuum and cryotemperatures on engineering .Materials: Biennial report', NASA

CR-77772, National Aeronautics and Space .Administration, George C. Marshall Space Flight Center. HuntsviUe, Alabama

(1966) (N66-35963 I 65 Keriin, E. E., Smith. E. T. "Measured effects of the various combinations of nuclear radiation, vacuum and cryotemperatures on engineering materials: Annual report', NAS.A

CR-58830. National Aeronautics and Space .Administration, George C. .Marshall Space Flight Center, HuntsviUe. Alabama

(1964) (N64-33043)

66 Keys. R. D., Kiefer. T. F.. Schwartzberg, F. R, Cryogenic

properties of high-strength glass-reinforced plastics'. Advances in Cryogen'C Engineering 11 (1966) 470

67 Krause, D. R. 'Development of lightweight material composites to insulate cryogenic tanks for 30-day storage

in outer space', NASA CR-1 23797, National Aeronautics and Space Administration, George C. Marshall Space Flight Center, HuntsviUe. Alabama ( 1972) (N72-30495)

68 Krause, D. R., Fredrickson, G. 0., Klevatt, P. L. Effects of

cychncal environments on high-performance multi-layer

insulation materials'. Society of Aerospace .Material and

Process Engineers National Technical Conference, Space

Shuttle .Materials, Vol 3 (1971) 639

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.. 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)

71 Levin, V. A,, Naumenkov, P. G., Shchitov, M. V. 'Some

properties of plastics at low temperatures'. /'/fliric/iejArw Massy 11 (1966) 64

72 Luikov, A. V., Vasiliev, L. L., 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 (1965) 314

73 Lyon, D. N., Parrish, W. R, 'Low temperature thermal conductivities of two high compressive strength materials'. Cryogenics 1 (1967) 21 74 Maher, L. E. 'Some problems arising from the use of hydrogen-fueUed propulsion systems'. / Roval Aeronautical Society 68 (1964) 765

75 McKannon. E, C, Gause, R, L. 'Effects of nuclear radiation and cryogenic temperatures on non-metaUic engineering

mzxsi\2t.\s\ J of Spacecraft 2 (1965) 558

76 Morris, E. E. 'Glass-fiber-reinforced metaUic tanks for cryogenic

service', 12th National Symposium., Society of Aerospace

Materials and Process Engineers, Vol 12 (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' yo/.Vcr 4 (1969) 970

78 Morris, E. E., Alfring, R- J. 'Cryogenic boron-filament-w-ound

pressure vessels', Composite Materials: Testing and Design, ASTM ST? 460, American Society for Testing and Materials (1969)430

79 Monis, E. E., Landes, R- E. 'Cryogenic glass-filament-wound tank evaluation', N.ASA CR-72948 (Final), National Aeronautics and Space Administration, NASA Lewis

Research Center, Cleveland, Ohio (1971) (N72-14696)

80 Mowers, R. E., Leib. J. H.. Sherman, S. 'Programme of testing nomnetaUic materials at cryogenic temperatures', Rocketdyne

Corporation Report R-3498, Rocket Propulsion Laboratories, Edwards, CaMfomia (1962) (.AD 294 772)

81 Nadler, M. .A., "V'oshino, S. Y.. Darms, F. J. 'Boron/epoxy

support strut for non-integral cryogenic tankage'. North

.American RockweU Space Division Report SD 68-99501

(1969) (see also 15th National Symposium S.AMPE AprU

1969 and North American RockweU Report SD 995 2, 1968) 82 Ndson, L. F. 'Compressive strength of Conolon 506 at + 75°F and - 320°F', Convair/Astronautics Report No 27E

1336 (1962) 83 Nelson, L. R, 'Mechanical properties of Adlock 851 at room

temperature, 1000°, -320° and -423°F', Convair/Astronautics Report No 55E 812 (1961) 84 Nelson, P. T., Archer, J. S. 'Graphite reinforced plastic EHF

antenna'. TRW Svsiems Group, Redondo Beach. California Report No 99906-7128-RO-11 (1969) 85 Patten, P. M. 'Internal insulation Uner alteration', Douglas

.Aircraft Company Report No S.M 45975 (1964) 86 Perkins-Elmer Optical Group. Norwalk, Connecticut, Workin-Progress on Contract No F33615-72-C-2033, Air Force Systems Command, Wright-Patterson Air Force Base, Ohio

87 Pink, E,, CampbeU, J. D. 'The effect of strain rate and temperature on the deformation behaviour of reinforced and unreinforced epoxy resin', O.xford University Department of

Eneineerins Report No 1040/72, Oxford, England (1972) (N73- 10568)

88 Pirgon, 0., VV'ostenholm, G. H.. Yates, B. "Thermal expansion

at elevated temperatures. IV. Carbon-tibre composite^ /

Physics D: AppI Phys 6 (19"73) 309

89 Pride, R. A.. Stein, B. A., Schmidt. F. W. 'Mechanical properties of polyimide-resin/glass-fiber laminates for various time, temperature and pressure exposures'. Technical Papers, 23rd .Annual Reiiiforced Plastics Technical and Management

Conference, Washington, DC (1968) section l7-c, 1

CRYOGENICS . DECEMBER 1975 37 719
90 Ratcliffe, E. H. 'Thermal conductivities of plastics with glass, asbestos and cellulasic fiber reinforcements', Appl Mat Res

5 (1966) 200

91 Roseland, L. M. 'Materials for cryogenic usage', Technical

Papers, 21st Annual Technical Conference, Society of the

Plastics Industry (1966) section 4-C, 1 92 Roseland, L. M. 'Investigation of structural properties at cryogenic temperatures of filament-wound pressure vessels containing both organic and glass filaments', Douglas Aircraft Corporation Report No SM^8409 (1966) 93 Ross, J. E. 'Fiberglass laminate - ultimate tensile and tlexural strength tests at room temperature, -100°F and 320°F', Convair Astronautics Report No 7E 1687 (1959) (AD 830 230) 94 Sanders, R. H., Weleff, W. 'Final report on GTR-17 effects of radiation on organic materials irradiated in liquid hydrogen'

Aerojet-General Corporation Report No RN-S-0317 (1967)

95 Sanger, M. J., Molho, R., Howard, W. W. 'Exploratory evaluation

of filament-wound composites for tankage of rocket oxidizers and fuels', AFML-TR-65-38 1

, Air Force Materials Laboratory,

Weright-Palterson Air Force Base, Ohio (1966) (AD 477 455) 96 Sanger, M. J., Reinhart, T. J. 'Development of filament- wound tankage for rocket oxidizers and fuels'. Technical

Papers, 12th National Symposium, Society of Aerospace

Material and Process Engineers (1967) section AS-7

97 Sewell, J. J., 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 98 Shriver, C. B. 'Design and fabrication of an internally insulated filament wound liquid hydrogen propellant tank', NASA

CR-1 27, National Aeronautics and Space Administration, Washington DC (1964) (N65-10775)

99 Soltysiak, D. J., Toth, J. M. 'Static fatigue of fiber glass pressure vessels from ambient to cryogenic temperatures'. Technical Papers, 22nd Annual Technical Conference,

Society of the Plastics Industry (1967) section 14-E

1 00 Speare, J. C. 'Preliminary sizing of filament-wound RNS

tanks'. Report No TOR-0066 (5759-07)-13, Space and

Missile Systems Organization, Air Force Systems Command,

Los Angeles Air Force Station, Los Angeles, California (1970) (AD 872 626)

101 Steinhauer, R. A. 'Linear thermal expansion of 828CL 181

cloth laminate', Douglas Aircraft Company Report No MP 1 1 979 (1961)

102 Stinnett, W. D. 'Cryogenic tensile properties of selected

materials', NASA CR-71751, AEC-NASA Space Nuclear

Propulsion Office, Report No 2712 (1964) (N66-22816)

103 Suezawa, Y., Hojo, H., Nakamura, K. 'Impact characteristics of fiberglass reinforced plastics at low temperatures', Kagaku

Kogaku (Chemical Engineering, Japan) 33 (1969) 1051 104 Toth, J. M. "Barrier films for filament-wound fiberglass cryogenic vessels', Advances in Cryogenic Engineering 1 (1964)537

1 05 Toth, J. M., Soltysiak, D. J. 'Investigation of smooth-bonded

metal liners for glass fiber filament-wound pressure vessels', NASA CR-72165 (Final), National Aeronautics and Space

Administration, Lewis Research Center, Cleveland, Ohio

(1967) (N67 -25070)

106 Toth, J. M., Sherman, W. C, Soltysiak, D. J. 'Investigation of smooth-bonded metal liners for glass fiber filament-wound

pressure vessels', Douglas Missile and Space Systems Division Report No SM49384, Quarterly Report No 3, Contract No NAS 3 6193, NASA Lewis Research Center, Cleveland, Ohio (1966)

1 07 Toth, J. M., Sherman, W. C, Soltysiak, D. J. 'Investigation of structural properties of fiber-glass filament-wound

pressure vessels at cryogenic temperatures', NASA

CR-54393, National Aeronautics and Space Administration, Lewis Research Center, Cleveland, Ohio (1965) (N65-35392)

108 Toth, J. M., Barber, J. R. 'Structural properties of glass-fiber filament-wound cryogenic pressure vessels', /Ic/vcnces Cryogenic Engineering 10 ( 1 965) 1 34 109 Voloshenko-Klimovitskii, Yu. Ya., Belyaev, Yu, A., L'vof,

B. S., Schpakovskaya, E. I. 'Strength of cold-hardening GRPs based on PN-1 resin under impact tension at normal

(20°C) and low (-196°C) temperatures', Plasticheski Massy

6 (1964) 39

1 10 Voloshenki-Klimovitskii, Yu. Ya., Belyaev, Yu. A., Korenkov,

Yu. A. 'Impact tensile tests on glass-fibre reinforced plastics

at normal and low temperatures', Plasticheski Massy No 5 (1963) 51

1 1 1 Watson, J. F., Christian, J. L., Hertz, J. Selection of

materials for cryogenic applications in missiles and aerospace

vehicles', Convair/ Astronautics Report No MRG 132-1

(1960) 112 Weleff, W. 'Effect of nuclear radiation and liquid hydrogen on mechanical properties of three phenolic materials', Advances in Cryogenic Engineering 1 1 (1966) 486

113 Weleff, W. 'Final report, GTR-16 radiation effects test on

structural materials at -423° F', Aerojet-General Corporation Report No RN-S-0290 (1966)

4. Handbooks and reviews 114 Coston, R. M. 'Handbook of thermal design data for multilayer

insulation systems', LMSC-A847882, Vol 2 (Final), George

C. Marshall Space Flight Center, Huntsvilie, Alabama (1967) (N67~34910)

1 15 Hertz, J. 'The effect of cryogenic temperatures on the mechanical properties of reinforced plastic laminates', Society oj Plastics Engineers J 2\ (1965) 181

1 16 Hertz, J., Knowles, D. 'Survey of thermal properties of

selected materials'. General Dynamics/Convair Report AAL-65-008 (AR-504-1-553) (1965) (N65-31775)

1 1 7 Jurevic, W. G., Rittenhouse, J. B. 'Structural applications handbook', AFML TR-67-332, Air Force Materials Laboratory,

Wright-Patterson Air Force Base, Ohio (1968) (AD 804 585)

1 1 8 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 (1965) (AD 469 126)

1 19 Maximovich, M., Scheck, W. G. 'Data summary and

reference file for graphite and boron reinforced composite

materials'. General Dynamics/Convair Report No GDCADBG71-006 (1971) (Contract NAS 8-26198, George C. Marshall Space Flight Center, Huntsvilie, Alabama)

1 20 Nored, D. L., Hennings. G., Sinclair, D. H., Smith, G. T., Smolak, G. R,, Stofan, A. J. 'Storage and handling of cryogenic fluids', NASA Special Publication SP-5053,

Proceedings of Conference on Selected Technology for the Petroleum Industry, Lewis Research Center, Cleveland, Ohio (1965) (N66-33674)

121 'Plastics for aerospace vehicles. Part 1. Reinforced plastics', M1L-HDBK-17A, Department of Defense, Washington, DC

(1971)

1 22 Rittenhouse, J. B., Singletary, J. B. 'Space materials handbook', 3rd edn, NASA Special Publication SP-3051, National Aeronautics and Space Administration, Washington, DC

(1969) (Limited publication as AF'ML-TR-68-205)

123 Sch wartzberg, F. R., Hertzog, R. G., Osgood, S. H. et al. 'Cryogenic materials data handbook (revised), Volume 2', AFML-TOR-64-280-Vol II (Revised), Air Force Materials Laboratory, Wright-Patterson Air Force Base, Ohio (1970) (AD 713 620)

5. Miscellaneous references

124 'Advanced composites design guide', 3rd edn, Vol 4: Materials, Air Force Materials Laboratory, Wright-Patterson Air Force

Base, Ohio (1973) 125 Kreider, K. G., Prewo, K. M. 'The transverse strength of boron

fibers'. Composite Materials, Testing and Design, ASTM

STP497 (1972) 539

1 26 Arp, v., Wilson, J. H., Windich, L., Sikora, P. 'Thermal

expansion of some engineering materials from 20 to 293 K', Cryogenics 2 (1962) 230

127 Novak, R- C, DeCrescente, M. D. 'Impact behaviour of

unidirectional resin matrix composites tested in the fiber direction'. Composite Materials, Testing-and-Design, ASTM

STP497 (1972) 311 128 Kastelic, J. R., Hiltner, A., Baer, E. 'Crazing, yielding and

fracture in polycarbonate and polyethelene terepthalate at low temperatures', J of Macro Molecular Science-Physics B7 (4) (1973) 679

720 38 CRYOGENICS . DECEMBER 1975

- 'i
Bibliography — property cross reference

Glass- Glass- Glass- Glass- Glass- Glass- Glass- Glass- Property Epoxy Polyester Phenolic Teflon Silicone Polyurethane Phenyl Silane PBI**

at" 1-1.3,2-2.2,1-1.3,2-2.2, 1-1,3,2-2.2, 2-2.2,37,67 1-1.3,2-2.2, 2-2.2,85 2-2.2,57,66, 2-2.2

3-3.5, 4, 5, 3-3.2, 3-4, 40, 54, 55, 57, 68, 80, 94, 24, 40, 57, 5.1,22,24, 3.5,34,40, 64-66,71,75, 102,113 64,66,68,97

28,33,35, 55,57,63, 83,93,97,

40,43,46, 64,66,75, 111,112

47,52,54, 109

55, 57, 66, 71,87

e\ 1-1.3,2-2.2,1-1.3,2-2.2, 1-1.3,2-2.2, 2-2.2,80 1-1.3,2-2.2, 2-2.2 2-2.2,57,66 2-2.2,24

3-3.5,4.5, 3,34,55,57, 54,55,57 24,57

5.1,24,28, 97,111

35, 55, 56, 57, 66, 87

E\ 1-1.3,2,3- 1-1.3,2,55, 1-1.3,55,57 1-1.3,2,57 2 2,57 2 3.5,28,55, 57

56, 57

e* 3-3.5,4,5, 3-3.5,64 64,65,112 80,89,102 64 85

5.1, 35,64

85

of" 1-1.3,2-2.2,1-1.3,2-2.2 1-1.3,2-2.2, 2-2.2,80 1-1.3,2-2.2, 2-2.2 2-2.2 2-2.2 3-3.5,4,40,3-3.3,3.5, 40,41,57,71, 40,57,97

57,66,71 40,57,66 83,93,97

e\ 1-1.3,2,3- 1-1.3,2,2.2, 1-1.3,2-2.2, 2-2.2,80 1-1.3,2-2.2 2-2.2 2-2.2 2-2.2

3.2,3.4, 57 57,83,93,97 57

3.5, 57 e[ 1-1.3,2 1-1.3,2 1-1.3,2 1-1.3,2 1-1.3,2 2 2

a<=" 1-1.3,2, 1-1.3,2,2.2, 1-1.3,2-2.2, 2-2.2,80 1-1.3,2-2.2, 2-2.2 2-2.2, 'ol 2-2.2,24

2.2, 3-3.2, 3-3.5, 34, 57 57, 82, 97 24, 57

3.4,3.5, 24, 57, 66

5= 1,1.2,1.3, 1-1.3,2,2.2, 1-1.3,2-2.2, 2-2.2,80 1-1.3,2-2.2, 2-2.2 2-2.2 2-2.2

2,2.2,57 57,66 57,82,97 57

a'' 3-3.2,3.4, 3-3.5 71 3.5,4,4.1,

5, 5.1, 22, 47, 71

T)* 1,1.3,2,2.1,1-1.3,2-2.2, 1-1.3,2-2.2 1-1,3,2 2 2 2

2.2, 66, 70, 34, 70

.by 99

2-2.2,3-3.2, 2 2 2 2

3.4, 3.5

a'* 4,71,103 62,109 71,97 80

X 3,4.1,14, 3,14,16-16.3,14,16-16.3, 14,16-16.3 14,16-16.3, 14,16-16.3 16,24

16-16.3,21, 90 49 23-25,38,

42,56,61,

63,90, 114 \UL 3,5,14,16, 3,14,16-16.3 14,16-16.3, 14,16-16.3,80 14,16-16.3, 5,85,91 14,16-16.3 16,16.2

16.2,16.3,28, 26,27,50, 67

69,85,91, 111 100,101,104,

114

Cp 15,16,16.2, 15,16-16.3 15,16-16.3, 15,16-16.3 15,16-16.3, 15,16-16.3 16-16.3,24

16.3,24,114 64 24

• includes fracture toughness

*• polybenzimidezole

CRYOGENICS . DECEMBER 1975 39 721
Bibliography—property cross reference (contd)

Property

GraphiteEpoxyBoron- EpoxyBoron- Aluminum

PRD-49

Epoxy

ctu 8, 9, 10, 10.2, 13,13.2,32,

47, 51, 58

13, 20, 22 47,53,58,

81

13, 13.3 29, 44,

59, 60

8, 9, 10, 10.1,

13, 51

13, 81 13, 13.3 44, 59, 60

13 20 13

8, 10, 10.1,

11, 12, 13, 13.2, 58

58

t' 8, 10.1 cu

a 1

3

13, 13.3 £C 13 13, 81 13, 13.3

8, 9, 10, 10.1,

10.2, 12, 13, 13.2, 47, 58

13, 22, 47,

58, 81

13, 13.3 59,60

t

rj 45 20

bv

13, 13.3

I* a' 9, 10.2 81 \ 13, 17, 18, 38 13, 38, 63, 81 AL/L 10, 13, 13.2,

17, 18, 36, 51, R9xJZ, 04OH, 00OR

13 13, 13.3 59,60

Cp 13 13 13, 13.3

* includes fracture toughness

Miscellaneous properties

Notch tensile strength Glass-Epoxy (3-3.2, 3.4, 3.5, 43,

46)

Vapour permeability Glass-Epoxy (28) Modulus of rigidity Glass-Epoxy (53), Glass-Teflon (80), Boron-Epoxy (53) Poissons ratio Glass-Epoxy (53), Boron-Epoxy (53) Proportional limit in tension Glass-Epoxy (55, 56), Glass- Polyester (1-1.3, 55, 56)

Static fatigue Glass-Epoxy (66, 99), PRD 49-Epoxy (44) Environmental effects G\ass-Epox\/ (4—4.4), GraphiteEpoxy (13, 13.2, 17, 58) Boron-Epoxy (13, 58), Boron-Aluminium (13, 36), PRD 49-Epoxy (44)

Electrical resistivity Graphite-Epoxy (13), Boron- Epoxy (13) Thermo-optical effects Graphite-Epoxy (13,84), Boron- Epoxy (13) Density Glass-Epoxy, Polyester, Phenolic, Silicone, Phenyl

Silane (14, 57), Glass-Teflon (14), Glass-Polybenzimidazole

(16), Graphite-Epoxy (9, 13), Graphite-Phenolic (14), Graphite-Polyimide (12.1) Radiation effects (13.2, 33, 37, 40, 51, 52, 64, 65, 75, 94,

112, 113)

Cryogen competability (33)

Miscellaneous composites

Glass-Polyimide a^" (24, 67, 68), a^" (89), a'' (12.1, 89)

AL/L (67)

, Glass-Melamine a*" (65), (65), X (90)

Glass- Viton a^" (68) Glass-Phenyl formaldenhyde a' (1 10), X (72), Cp (72) Si02-Epoxy a'" (31), AL/L (28) Graphite-Aluminium a*" (thermal cycling effects 6—6.2)

Graphite-Polyimide a (14, 16-16.3) AL/L (14, 16-16.3)

Steel-Aluminium a*" (7, 7.1), e' (7.1), a* (7, 7.1) Steel-Epoxy a'" (2, 30), (2, 30), a^" (2), a^^" (2), (2),nt (2) Boron/Steel-Aluminium\ a'", f*, ct'=", £^ a^^ , AL/L/C^

Boron/Titanium- /(1 3, 13.3), a»', e«, a7M13)

Aluroininium ) Potassium titanate-epoxy X (16), Cp (16)

Pressure vessel applications

Glass- filament (19, 20, 39, 42, 46-48, 53, 67, 69, 76, 77, 91, 92, 95, 96, 98-100, 104-108, 113)

Graphite-filament (8, 10, 10.1, 45, 47, 100)

Boron-filament (20, 47, 53, 78) PRD 40-filament (59, 60)

722

40 CRYOGENICS. DECEMBER 1975
SEMI-ANNUAL REPORT ON MATERIALS RESEARCH

IN SUPPORT OF SUPERCONDUCTING MACHINERY

ELASTIC PROPERTIES OF ENGINEERING MATERIALS

AT CRYOGENIC TEMPERATURES

H. M. Ledbetter and D. T. Read

Cryogenics Division

Institute for Basic Standards

National Bureau of Standards

Boulder, Colorado 80302

April 1976

41
Suimnary; Elastic Properties

During the six months preceding March 1976, the following studies were completed:

(1) 2000-series aluminum alloys . Two alloys -- 2014 (Al-Cu-Si-Mg) and

2219 (Al-Cu) were studied in their precipitation-hardened conditions by a pulse-echo method between room temperature and liquid-helium temperature.

These alloys exhibit similar elastic properties and regular behavior with

respect to temperature. They are slightly stiff er elastically than some of

the other aluminum alloys that were reported on previously in these reports.

Results of the study are given in an accompanying manuscript "Low-temperature

elastic properties of aluminum alloys 2014 and 2019" by D. T. Read and H. M.

Ledbetter.

(2) Face-centered-cubic , iron-nickel, high-permeability alloy . Iron47.5 nickel was studied between room temperature and liquid-helium temperature by a pulse-echo method. Besides invar, iron-36 nickel, this is the

only material of this type for which low-temperature elastic constants have

been determined. Contrary to some previous reports, the room-temperature

thermo-elastic coefficients are normal for this material. Results of the

study are given in an accompanying manuscript "Low-temperature elastic properties of iron-47.5 nickel" by H. M. Ledbetter and D. T. Read.

(3) Nickel-chromium-iron-molybdenum alloy . There is an accompanying

reprint of our previously described results on Inconel 718: W. F. Weston

and H. M. Ledbetter, "Low- temperature elastic properties of a nickelchromium-iron-molybdenum alloy," Mater. Sci. Engg. 2_0 (1975) 287-90.

(4) Aluminum alloys 1100, 5083, 7005, 7075 . There is an accompanying

reprint of our previously described results on four aluminum alloys: E. R.

Naimon, H. M. Ledbetter, and W. F. Weston, "Low-temperature elastic properties of four wrought and annealed aluminium alloys," J. Mater. Sci. 1_0 (1975)

1309-16.

(5) Copper-nickel alloys . Our study on copper, copper-10 nickel, and

copper-30 nickel was published: H. M. Ledbetter and W. F. Weston, "Lowtemperature elastic properties of some copper-nickel alloys," 1975 Ultrasonics

Symp. Proc, IEEE Cat. No. 75 CHO 994-45U, 623-7. A reprint is not included

here because it is indentical to the preprint version given previously.

(6) Superconducting-coil composite . There is an accompanying reprint

of our previously described results on a niobium-titanium, copper-stabilized,

epoxy-impregnated , superconducting-coil composite: W. F. Weston, "Lowtemperature elastic constants of a superconducting coil composite," J. Appl

.

Phys. 4_6 (1975) 4458-65.

42
LOW-TEi^PERATURE ELASTIC CONSTANTS OF PRECIPITATIONHARDENED ALUMINUM ALLOYS 2014 and 2219"*"

D. T. Read* and H. M. Ledbetter

Cryogenics Division

Institute for Basic Standards

National Bureau of Standards

Boulder, Colorado 80302

+ Contribution of NBS, not subject to copyright. This work was partially supported by the Advanced Research Projects Agency

of the U. S. Dept. of Defense.

* NRC-NBS Postdoctoral Research Associate, 1975-76.

A3
Abstract

Elastic properties of aluminum alloys 2014 and 2219 were studied

in their precipitation-hardened conditions between 4 and 300 K using

an ultrasonic pulse-echo superposition technique. Results are given

for longitudinal sound velocity, transverse sound velocity. Young's

modulus, shear modulus, bulk modulus, and Poisson's ratio. The Young's

moduli increase about eleven percent on cooling. The shear moduli are

about ten percent higher than for unalloyed aluminum. All the elastic

constants show regular temperature behavior.

Key words: Aluminum alloys; bulk modulus; compressibility; Debye

temperature; Poisson's ratio; precipitation-hardening alloys; shear

modulus; sound velocity; Young's modulus.

44
Introduction

The combination of a high ratio of strength to weight and high

toughness at low temperatures makes many aluminum alloys attractive for

cryogenic structural applications. In this report, the cryogenic elastic properties of two aluminum alloys, 2014 and 2219, used frequently

at low temperatures, are described. All 2000-series alloys are precipitation hardenable, containing copper as the principal alloying element.

Alloy 2014 has better mechanical properties, but alloy 2219 has

superior weldability. This is believed to be the first report of a

complete set of low-temperature polycrystal 1 ine elastic constants for a

series of 2000-type aluminum alloys.

Elastic constants are necessary parameters in the design of critical load-bearing members; they relate the stress applied to an object

to its change in dimensions produced by the stress. Elastic constants

are also used in analyzing the fracture properties of a material. The

factor E/(l - v ), where E is Young's modulus and v is Poisson's ratio,

occurs frequently when the strength properties of materials are considered.

Some elastic-property data on these alloys were reported previously[l].

However, the inaccuracies of those data were high, and the shear modulus

was not reported. Thus, neither the bulk modulus nor Poisson's ratio,

which are frequently of interest in engineering design, can be computed.

In the present study, an ultrasonic (10 MHz) pulse-superposition

method[2] was used for determining the velocity v of an ultrasonic pulse

propagated through the specimen. Appropriate moduli C were calculated

2 using equations of the type C = pv , where p is the mass density. This

indirect technique is more accurate than the more familiar direct

45
measurement of the change in length of a specimen under load where large

errors can arise from mechanical misalignments. Ultrasonic methods have

many advantages: small specimens are sufficient, so that 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;

5 relative imprecision is low, about one part in 10 for the velocities;

laboratory-to-laboratory variations of the elastic constants are typically a few percent or less; and the relevant properties of the quartz

transducer used to generate and detect ultrasonic pulses are affected

only slightly by temperature.

Materials

Aluminum alloys 2014 and 2219 are precipi tation-hardenable alloys

containing copper and other alloy elements. Their chemical compositions

are given in Table 1. The 2014 alloy was tested in the T652 condition:

solution heat-treated, artificially aged, and stress relieved by rolling. The 2219 alloy was tested in the T87 condition: solution heat

treated, cold worked, and artificially aged. The 2014-T652 and 2219-

T87 alloys had Rockwell hardnesses of B 80.6 and B 81.1 and mass densities of 2.810 and 2.835 g/cm^, respectively. Metal lographic examination

of these alloys revealed severely elongated grains. Surfaces perpendicular to the rolling direction showed nearly equi-axed grains with diameters of about 0.02 cm in the 2014 alloy and 0.03 cm in the 2219 alloy.

Surfaces parallel to the rolling direction showed grains so severely

elongated that their dimensions could not be determined accurately.

Procedures

The specimen materials were obtained from commercial sources; the

46
2014 alloy was obtained as a 7.6-cm thick forging, and the 2219 alloy

was received in the form of a 3.8-cm thick plate. The hardnesses were

determined by a standard technique; the mass densities were determined

by the method of Archimedes using distilled water as a standard.

Quartz transducers were bonded to the specimens with phenyl salicylate

(salol) for room temperature measurements and with stopcock grease for

the initial part of the low-temperature measurements. These bonds failed

around 70 K, so that a silicone fluid (viscosity = 20 kPa-s at 25° C) was

used for bonding at very low temperatures. The low-temperature apparatus

was described previously[3]

.

Elastic constants were determined from sound-velocity measurements

between room temperature and liquid-helium temperature. The inaccuracy of

the velocities is estimated to be less than + 1%.

Transducer and bond effects were minimized by measuring the transit

times for both the one-transducer (usual) case and the two-transducer case,

using "identical" transducers and bonds. The corrections were about one

percent.

Low-temperature data were obtained by a pulse-echo-superposition

method[2], which gives the ratio of the low-temperature velocity to the

room-temperature velocity with an imprecision of about 0.001%.

The main limitation on the usefulness of the present results is the

batch-to-batch variation of commercial materials. Due to this factor,

variations in the elastic constants as large as one percent could be expected,

but temperature coefficients should be unaffected. Thus, measurements on

similar alloys at room temperature should establish their elastic constants

over the entire 0-300 K region.

47
Results and Discussion

The measured longitudinal moduli C = pv^ for both alloys are shown

in Fig. 1. The small steps between adjacent data points for the 2219

alloy at 80 K and 140 K are not representative of the material's elastic behavior, but are artifacts of the measurement process. They are

due to the use of two different bonding agents to seal the quartz transducer to the specimen for the low temperature measurements. The temperature dependences of the longitudinal moduli of both the 2014 and the

2219 alloys are quite regular.

The measured transverse moauli are shown in Fig. 2. Again, the

small step at 120 K in the modulus of the 2219 alloy is a measurement

artifact. The temperature dependences of the transverse moduli of both

alloys are also quite regular. The data shown in Figs. 1 and 2 were

fitted to a function of temperature suggested by Varshni[4]:

C = CO - i—s—

exp(t/T)-r

where C is any elastic stiffness constant (C-| and in this case), C^,

s, and t are the fitting parameters, and T is the temperature. The fitted

curves are shown along with the data in Figs. 1 and 2. Values of the fitting parameters are given in Table 2, Average differences between the mea

sured moduli and the fitted curves were about 0.05%.

The elastic De^-ye temperatures for these alloys were calculated

from their elastic constants at T = 0 K, and they are given in Table 2.

48
As expected from their higher shear moduli (discussed below), the alloys

have higher Debye temperatures than "pure" aluminum where £ = 431 K.

Values of the longitudinal and transverse moduli calculated using

the fitted Varshni parameters were used to calculate other elastic constants. The additional elastic constants reported here: the shear modulus

G, the Young's modulus E, the bulk modulus B, and the Poisson's ratio v,

are given by the following formulas [5]:

G = = Qi\ , (2)

E = 3G(C^ - 4/3 C^)/(C^ - C^), (3)

B = - 4/3 C^, (4)

and

V = 1/2(C-, - 2 C^)/(C^ - C^). (5)

The four elastic constants obtained using these relations are shown as

functions of temperature in Figs. 2-5. Values of these elastic constants

at selected temperatures are given in Table 3, along with some previous

results.

As shown by the data in Table 3, changes in the elastic constants of

aluminum alloys 2014 and 2219 between 300 and 0 K are about 12"'^ for E and

G, and about 4"; for B and v. These changes are much larger than those

49
observed in alloys based on iron[6] and copper[7] , two other common base

metals. Most of the changes in the elastic constants of these aluminum

alloys with temperature occur above liquid-nitrogen temperature, 77 K.

The outstanding feature of the data shown in Figs. 1-5 is the

regular temperature dependence of all the elastic constants. These two

aluminum alloys are good examples of materials with simple, ordinary

elastic behavior with respect to temperature. This regular behavior was

not evident from some previous reports on these alloys.

The bulk moduli of the 2014 and 2219 alloys differ by 0.3% and 3.9%

from that of pure aluminum, while the shear moduli of these alloys differ

by 8.5% and 11.9% from that of pure aluminum[8]. Thus, the alloying

elements increase not only the yield strength but also the shear modulus

of the base material. Alloying effects in these materials are complicated

by the large number of alloying elements, but it is clear that the alloying raises the shear moduli much more than is expected from a simple

rule of mixtures, while the bulk modulus is affected only slightly.

These alloys were not examined for anisotropic elastic behavior

because aluminum single crystals are only slightly ani sotropic[8] . Even

strongly textured polycrystal 1 ine aggregates of aluminum should exhibit

nearly isotropic elastic properties.

Compared to other aluminum alloys that were studied previously[9]

,

2000-series alloys have higher longitudinal and transverse moduli, resulting in higher Young's moduli and higher bulk moduli than the other alloys;

but they have the same Poisson's ratio. The higher elastic moduli of the

2000-series alloys may be useful in critical design situations.

Conclusions

(1) With respect to unalloyed aluminum, alloys 2014 and 2019 have

50
higher Young's moduli ('v 9?0 > higher shear moduli 10%), higher bulk

moduli 2%), and lower Poisson ratios ('v 3%).

(2) These elastic moduli are also higher than those reported previously for 5000-series and 7000-series aluminum alloys.

(3) Alloy 2219 has a slightly higher shear modulus, but a slightly

lower bulk modulus, than alloy 2014.

(5) The temperature behaviors of the two alloys are very similar.

Between 300 and 0 K, B increases about 4%, E increases about 11%, G

increases about 12%, and v decreases about 3 %.

51
References

1. Christian, J. L. and Watson, J. F. , "Mechanical Properties of

Several 2000 and 6000 Series Aluminum Alloys," Advan. Cryo. Eng .,

Vol. 10, 1965, pp. 63-76.

2. McSkimin, H. J., "Pulse Superposition Method for Measuring Ultrasonic

Wave Velocities in Solids," J. Acoust. Soc. Amer . , Vol. 33, 1961,

pp. 12-16.

3. Naimon, E. R. , Weston, W. F. , and Ledbetter, H. M. , "Elastic Properties of Two Titanium Alloys at Low Temperatures," Cryogenics ,

Vol. 14, 1974, pp. 246-249.

4. Varshni, Y. P., "Temperature Dependence of the Elastic Constants,"

Phys. Rev. B , Vol. 2, pp. 3952-3958.

5. Landau, L. D. and Lifshitz, E. M. , " Theory of Elasticity ," Pergamon,

London, 1959, pp. 13, 99.

6. Ledbetter, H. M. and Reed, R. P., "Elastic Properties of Metals and

Alloys. I. Iron, Nickel, and Iron-Nickel Alloys," J. Phys. Chem .

Ref. Data , Vol. 2, 1973, pp. 531-618.

"/. Ledbetter, H. M. and Naimon, E. R., "Elastic Properties of Metals

and Alloys. II. Copper," J. Phys. Chem. Ref. Data , Vol. 3, 1974,

pp. 897-935.

8. Kamm, G. N. and Alers, G. A., "Low-Temperature Elastic Moduli of

Aluminum," J. Appl . Phys . , Vol. 35, 1964, pp. 327-330. Polycrystalline averages were calculated from the single-crystal data in this

reference using the Voigt-Reuss-Hill-arithmetic method.

9. Naimon, E. R. , Ledbetter, H. M., and Weston, W. F., "Low-Temperature

Elastic Properties of Four Wrought and Annealed Aluminium Alloys,"

J. Mat. Sci. . Vol. 10, 1975, pp. 1309-1316.

52
Table 1 Chemical compositions of studied

aluminum alloys, wt. %^

Al loy Cu Fe Mg Mn Sn Ti V Zn Zr Al

2014 4.4 0.05 0.8 0.8 Bal

2219 6.4 0.20 0.01 0.26 0.15 0.16 0.12 0.09 0.16 Bal

3 For 2014, nominal composition. For 2219, plasma arc analysis.

Table 2 Varshni-function parameters determined from a least-squares

fit of the data (see equation 1) and elastic Debye temperature, Q

Alloy Mode C°(10^^ Nm-2) s(10^^ Nm"2) t(K) G(K)

2014 ' 1 1 .241 0.0781 217 439

t 0.3152 0.0339 212

2219 1 1.220 0.0859 230 443

t 0.3250 0.0345 209

53
at selected temperatures in units of Nm"'^, except

V, which is dimensionless

Previous

Present Results results^

Alloy Temperature (K) B E G V E

2014 300 .791 .757 .282 .341 .73

" 200 .805 .794 .297 .336 .74

100 .817 .827 .311 .331 .76

0 .821 .838 .315 .330 —

2219 300 .758 .774 . .291 .330 .72

200 .772 .812 .306 .325 .72

100 .784 .845 .320 .320 .74

0 .787 .857 .325 .318

^ 2014 alloy in the T-6 condition, average of longitudinal and transverse, interpolated data; 2219 alloy in the T-81 condition, average

of longitudinal and transverse, interpolated data.

List of Figures

Fig. 1. Longitudinal modulus C-j = pv^ for two 2000-series aluminum

alloys in precipitation-hardened condition as a function of

temperature.

Fig. 2. Shear modulus G = = pv^ for two aluminum alloys.

Fig. 3. Young's modulus for two aluminum alloys.

Fig. 4. Bulk modulus for two aluminum alloys.

Fig. 5. Poisson's ratio for two aluminum alloys.

54
1.26 1 1 r r

55
0.74

0 50 100 150 200 250 300

TEMPERATURE (K)

Fig. 4. Bulk modulus for two aluminum alloys.

36

Low-Temperature Elastic Properties of an Iron-47.5 Nickel Alloy"*"

H. M. Ledbetter and D. T. Read*

Cryogenics Division

Institute for Basic Standards

National Bureau of Standards

Boulder, Colorado 80302

ABSTRACT

Elastic properties of a polycrystalline iron-47.5 wt. pet. nickel alloy were studied between room temperature and liquid-helium temperature. Roomtemperature longitudinal and transverse ultrasonic velocities were determined by a pulse-echo technique. Low-temperature velocities were determined

by a pulse-echo-superposition technique. The reported elastic constants are longitudinal modulus, shear modulus, Young's modulus, bulk modulus, Poisson

ratio, and elastic Debye temperature. Low-temperature elastic constants of

face-centered cubic iron-nickel alloys are reviewed. Contrary to some previous reports, the elastic constants behave regularly in the temperature

range studied.

Key words: Bulk modulus; compressibility; Debye temperature; iron alloy;

nickel alloy; Poisson ratio; shear modulus; sound velocity; Young's modulus,

+ Contribution of NBS , not subject to copyright.

^ NRC-NBS Postdoctoral Research Associate, 1975-6.

58
INTRODUCTION

Elastic properties of an iron-47.5 nickel alloy were studied between

room temperature and liquid-helium temperature. Elastic constants were determined by measuring the velocities of both longitudinal and transverse

ultrasonic waves in a polycrystalline aggregate. The following constants

are reported here: = longitudinal modulus, G = shear modulus, B = bulk

modulus, E = Young's modulus, v = Poisson's ratio, and 0 = elastic Debye

temperature

.

Iron alloys containing 30-100 percent nickel have a face-centered cubic

crystal structure, and they are magnetic at room temperature. These alloys

have many interesting properties that are exploited technologically . Many

of these properties involve magnetostriction, the coupling between magnetization and elastic strain. Recently, some of the alloys, developed originally

for high-temperature applications, have been used cryogenically . The alloy

reported on here--Fe-47 . 5Ni--is a candidate material for shaping magnetic

fields in superconducting machinery. Desirable properties for such an application include: (a) f erromagnetism with high permeability, (b) absence of

phase transitions, (c) at least moderate yield strength, and (d) adequate

low-temperature mechanical properties, including both fatigue and fracture.

Elastic constants are interesting for two principal reasons. First,

elastic constants are related directly to interatom.ic forces; thus, they are connected with a variety of solid-state pnenomena, including maximum attainable strengths, phase stabilities, and lattice specific heats. Second,

elastic constants are essential design parameters; to compute deflections

due to applied loads or stresses due to temperature changes of constrained

components, the elastic constants must be known.

While the elastic constants of iron-nickel alloys have been studied

extensively,*^ no low-temperature elastic data exist for any of the facecentered cubic alloys except invar, 5 iron-36 nickel. Because of the many

elastic anomalies exhibited by iron-nickel alloys,^ characteristic of magnetic materials, all predictive schemes for elastic constants are unreliable

for Fe-47.5Ni. Therefore, detailed experimental studies are required in this

case.

EXPERIMENTAL

Material

The studied material was obtained from a commercial source in the form

of a 10 X 23 X 108 cm ingot. The chemical (mill) analysis of the material

by weight was: 0.012C, 0.33Mn, 0.25Si, 0.002P, 0.004S, and 47.47Ni. On an atomic basis, the alloy contained 46.2 Ni. The grain size was 0.12 mm determined by an intercept method. The diamond-pyramid hardness number was 12 3 for a 1 kg load. The mass density was 8.19 g/cm^ , as determined by Archimedes '

s

method using distilled water as a standard.

Room-Temperature Sound Velocities

Room-temperature sound velocities were measured by a pulse-echo technique. 7 Briefly, a quartz piezoelectric transducer with a fundamental

resonance frequency of 10 MHz was cemented with phenyl salycilate (salol)

to one end of a specimen having flat and parallel faces. The specimen in

this case was a 1.25 X 1.25 X 1.6 cm parallepiped with opposite faces ground

flat and parallel within 2.5 um. Ultrasonic pulses about one ps in duration

were sent into the specimen by electrically exciting the transducer. The

pulses propagated through the specimen, reflected from the end, and propagated

59
back. The echoes were detected by the transducer and displayed on an oscilloscope. The sound velocity was computed by

V = 2£/t (1)

where £ is the specimen length and t is the transit (round-trip) time. On

the oscilloscope, t was the time between subsequent echoes. The oscilloscope

was calibrated against a precision time-mark generator. An x-cut quartz transducer was used for longitudinal waves, and an ac-cut quartz transducer was used for transverse waves. Low-Temperature Sound Velocities

Low-temperature measurements of the sound velocities were made by a pulse-echo-superposition technique.^ Briefly, the repetition rate of the

pulse was increased so that each pulse coincided with the second echo of the

preceding pulse. Since the excitation voltage was large compared with the

echo voltages, the oscilloscope display consisted of alternating pulses of

excitation voltages and "echo" voltages where the "echo" voltage represented

the sum of all odd-numbered echoes of the non-superimposed case. Because of

interference effects, the envelope of the summed odd-numbered echoes is highly

sensitive to small changes, cuased in this case by cooling, in the ultrasonic

velocity

.

The transducer-specimen bonding material was a stopcock grease. Temperatures were monitored with a chromel-constantan thermocouple placed near the

specimen. Cooling rates were about 2 K/min. The specimen holder was described

previously. 9 Cooling was achieved by lowering the specimen-holder assembly

stepwise into the ullage of a helium dewar . No thermal contraction corrections were made. For this alloy, the maximum thermal contraction correction

to the elastic constant, which applied at T = 0 K, is 0.3 percent. No cor- rection was made for the transducer-cement-coupling phase shift; the McSkimin^O

analysis gives a correction of less than 0.5 percent in the velocity, assuming

a maximum phase shift of it. The transit time was corrected to allow for the

thickness of the transducer; this correction is approximately one cycle at

10 MHz; thus, the observed longitudinal and transverse wave transit times

were reduced about 1.0 and 0.5 percents, respectively.

Room-temperature sound velocities and elastic constants are given in

Table 1. Both velocities were measured in three orthogonal directions and

averaged to allow for texture. For each direction, shear velocities were measured for two orthogonal polarizations.

Longitudinal and shear elastic constants were computed from the velocities according to

RESULTS

C£ = PV£ (2)

and

G = (3)

where p is the mass density.

60
The temperature variation of these is shown in Figs. 1 and 2. Curves in

those figures are least-squares fits of the data to the Varshni^l relationship:

^ = ^° - exp(t/T)-l

where C is C; or C^, 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 and Einstein oscillator model, it can be shown that t is the Einstein temperature. Parameters

C°, s, and t are given in Table 2. Average differences between measured and curve values are 0.04 and 0.05 percent for the longitudinal and transverse

cases, respectively.

Temperature variations of Young's modulus, the bulk modulus, and Poisson's

ratio are shown in Figs. 3-5. These elastic constants were computed from

Eq. (4) and the parameters in Table 2 using the relationships:

B = C- - 4/3G, (5)

E = 9GB/ (G + 3B) , (6)

and

V = (E/2G)-1. (7)

The elastic Debye temperature was computed according to

Q = h V (8) K 4 -A ^vg

where h is Planck's constant, k is Boltzmann's constant, A is the effective

atomic weight (57.00), and the average sound velocity is given by

v-3 + 2v-3 -1/

'avg = ^-^) ^

(9)

The result was found to be 3 = -i32 k, using the elastic data extrapolated to

T = 0 K.

DISCUSSION

A. Temperature Dependence of the Elastic Constants

In the temperature region studied, T = 0 - 300 K, all the polycrystalline

elastic constants behave regularly with respect to temperature changes. The

longitudinal, Young's, shear, and bulk moduli all increase monotonically with

decreasing temperature, show linear behavior at higher temperatures, show a continuously decreasing slope with decreasing temperature, approach relative

flatness at low temperatures, and approach zero slope at T = 0 K, in accordance with the third law of thermodynamics. The Poisson ratio (which can be

considered to be a function of the ratio of any combination of Cj, B, E, or

G) shows the same behavior, except that it decreases with decreasing temperature, which is usual for Poisson's ratio.

61
The temperature dependences of Young's modulus and the shear modulus

are very similar. This can be understood simply as follows. Differentiation

of Eq. (8) with the approximation that v - 1/3 gives:

E ' dT ~ 9 G

• dT 9 ' B ' dT

Thus, for usual values of the Poisson ratio, changes in Young's modulus

depend almost entirely in changes in the shear modulus and only slightly on changes in the bulk modulus.

At higher temperatures, as the Curie temperature of the alloy (T^ ~ 730 K)

is approached, the temperature behavior of the elastic constants may become

irregular.

B. Other Topics

In a final manuscript, several additional topics will be discussed.

These include: (i) a comparison of the present temperature coefficients with

those reported previously by other researchers; (ii) a review of the lowtemperature elastic-constant data for f ace-centered-cubic iron-nickel alloys;

(iii) a brief discussion of the elastic Debye temperature of fee iron-nickel

alloys; (iv) the relationship between single-crystal and polycrystal elastic

constants for the case of magnetic materials; and (v) the effect of magnetism

and magnetic fields on the elastic constants per se .

CONCLUSION

Contrary to some previous reports, all of the polycrystalline elastic

constants of iron-47.5 nickel exhibit regular behavior with respect to temperature^ in the region T = 0 - 300 K.

ACKNOWLEDGMENT

This work was supported in part by the Advanced Research Projects Agency

of the U.S. Department of Defense.

REFERENCES

1. S. J. Rosenberg: Nickel and Its Alloys , National Bureau of Standards

Monograph 106, May 196 8.

2. R. M. Bozorth: Ferromagnet ism , pp. 102-28, Van Nostrand, New York, 1951.

3. S. Chickazumi: Physics of Magnetism , pp. 495-97, Wiley, New York, 1964.

4. H. M. Ledbetter and R. P. Reed: J. Phys. Chem. Ref. Data , 19 74, Vol. 2,

pp. 531-618.

5. H. M. Ledbetter, E. R. Naimon, and W. F. Weston: Proc. ICMC-CEC (Kingston,

Ontario, July 1975), forthcoming.

6. G. Hausch and H. Warlimont: Acta Metall. , 1973, Vol. 21, pp. 401-14.

62
8.

9.

R. T. Beyer and S. V. Letcher: Physical Ultrasonics, pp. 79-87, Academic,

New York, 1969.

H, McSkimin: J. Acoust. Soc. Amer.

,

1961, Vol. 33, pp. 12-16,

E. R. Naimon, W. F. Weston, and H. M. Ledbetter: Cryogenics , 1974, Vol. 14,

pp. 246-49.

Table I. Room-temperature Acoustic-wave

Velocities in Fe-47.5 Ni, 10^ cm/s

Propagation

Direction

X

y

z

"1

0.5673

0.5648

0.5589

""t^1

0.2882

0.2881

0.2848

"-2

0.2879

0.2865

0.2841

Table II. Parameters Determined from

the Varshni Equation

C°, 10" N/m^

S, 10" N/m

t , K

Longitudinal

2.654

0.1036

221. 9

Transverse

0.7047

0.05959

264 . 4

LIST OF FIGURES

Fig. 1. Longitudinal modulus C]_ = pv^ for iron-47.5 nickel as a function

of temperature.

2 Fig. 2. Shear modulus G = = pv^ for iron 47.5 nickel.

Fig. 3. Young's modulus for iron-47.5 nickel.

Fig. 4. Bulk modulus for iron-47.5 nickel.

Fig. 5. Poisson's ratio for iron-47.5 nickel.

63
64
1.88 I

1 1 1 1 r

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

TEMPERATURE (K)

Fig. 3. Young's TOdulus for iron-47.5 nickel.

TEMPERATURE (K)

i

Fig. 4. Bulk modulus for iron-47.5 nickel.

0.328

50 100 150 200 250 300

TEMPERATURE (K)

Fig. 5. Poisson's ratio for iron-47.5 nickel.

65
Materials Science and Engineering, 20 (1975) 287—290

© Elsevier Sequoia S.A., Lausanne — Printed in The Netherlands

Short Communication

Low-temperature Elastic Properties of a Nickel—Chromium—Iron—Molybdenum Alloy*

W.F. WESTON+ and H.M. LEDBETTER

Cryogenics Division, Institute for Basic Standards, National Bureau of Standards, Boulder, Colorado 80302 (U.S.A.)

(Received in revised form May 12, 1975)

The low-temperature elastic properties of two

nickel - chromium - iron alloys, Inconel 600**

and Inconel X-750, were reported recently [1].

The same properties — the longitudinal modulus,

Young's modulus, the shear modulus, the bulk

modulus (reciprocal compressibility) and

Poisson's ratio — are reported here for a nickel - chromium - iron - molybdenum alloy, Inconel

718. Except for being slightly softer elastically,

it resembles Inconel X-750.

The low-temperature elastic properties of

Inconel 718 are of interest mainly because it is a candidate material for cryogenic structural

apphcations [2]. The elastic properties are important both for the design of structural

components and for understanding the basic

physical properties of a material. From the

temperature-dependent elastic constants, deflections can be predicted for any combination

* Contribution of NBS, not subject to copyright.

NRC-NBS Postdoctoral Research Associate,

1973 - 74. ** Tradenames are used to identify the materials

tested; they are not NBS endorsements of particular products.

of stress and temperature. Also, the Debye

characteristic temperature can be calculated

from the elastic constants. The Debye temperature relates in turn to a wide vsiriety of solid- state phenomena [3].

Inconel 718 is made precipitation hardenable by its niobium-plus-tantalum content.

(Inconel X-750 is made precipitation hardenable by its aluminum-plus-titanium content.)

It differs from the nickel - chromium - iron

alloys in having superior mechanical properties

such as yield, creep, rupture and fatigue

strengths [2]. Its sluggish response to precipitation hardening permits annealing and welding

without significant property changes.

EXPERIMENTAL METHODS

Ultrasonic (10 MHz) longitudinal and trans- verse sound-wave velocities were measured

between room temperature and liquid-helium

temperature (4 K) using a pulse-echo -super- position method [4]. Experimental procedures were identical with those reported previously

[1] except that the material was obtained from

TABLE 1 Composition and properties of the alloy

Chemical composition, mill analysis Hardness Mass density at ,wt.%) (DPH No., 1 kg load) 294 K(g/cm3)

Ni Cr Fe Nb + Ta Mo Ti 425 8.229

53.73 ^ 18.49 17.62 5.17 2.98 1.01

Al C Mn Si S Cu Condition

:

As-received; hot-rolled 0.58 0.05 0.08 0.17 0.007 0.04

66
Z6S

a commercial source in the form of 3/4-inch

(1.9 -cm) rods. Chemical and physical data on

the material are given in Table 1.

RESULTS

The longitudinal modulus

Ci=pvf (1)

is shown versus temperature in Fig. 1, where p

is the mass density and l'i is the longitudinal

sound-wave velocity. The transverse modulus

C,=pv^ = G (2)

is shown versus temperature in Fig. 2, where

t't is the transverse sound-wave velocity and G

is the shear modulus. Young's modulus E, the

bulk modulus B and Poisson's ratio p were

calculated from the formulas [5]

:

i: = 3Ct(Ci-tCJ/(Ci-Ct), (3)

B = C,-iC„ (4)

u =(l/2)(Ci-2Ct)/(Ci-Ct); (5)

and these elastic constants are shown versus temperature in Figs. 3 - 5. For comparison,

Figs. 1 - 5 also show as dashed lines the elastic constants of Inconel X-750, which were re- ported previously [1].

The temperature dependences of both Q

and Ct were fitted to a theoretical relationship

suggested by Varshni [6]:

C = C° -s/{e"'^-l) (6)

where C°, s, and t are adjustable parameters

Fig. 2. Transverse (= shear) modulus vs. temperature

for two Inconel alloys.

S ?.050 -

= ?.ooo -

"1 1 1 r

S '-'50

J L J L

0 so 100 150 200 2S0 300 TEMPERATURE (K) Fig. 3. Young's modulus vs. temperature for two

Inconel alloys.

^ .' -7tQ —

Fig. 1. Longitudinal modulus us. temperature for two

Inconel alloys,

SCO 250 300

Fig. 4. Bulk modulus (reciprocal compressibility) lis. temperature for two Inconel alloys.

67
289

0- 30S -

TEMPf fiSIllBE 1«)

Fig. 5. Poisson's ratio vs. temperature for two Inconel

alloys.

and T is temperature. The average difference

between curve values and measured values was 0.06 and 0.05% for the transverse and the

longitudinal modes, respectively.

The elastic Debye temperature was calculated from C) and Ct using standard formulas

[1] and was found to be 0 = 446 K at absolute

zero. For comparison, 6 (nickel) = 476 K and

6 (Inconel X-750) = 474 K.

DISCUSSION

Basically, the low-temperature elastic properties of Inconel 718 are very similar to those

of Inconel X-750. The higher-iron, lower-nickel

content of Inconel 718 accounts for its slightly

lower elastic Debye temperature. Inconel 718

is also similar to Inconel X-750 in the temperature dependences of its elastic constants for

the cases of E, G, and Q, which show regular

behavior. By regular behavior is meant: con- tinuously decreasing with increasing temperature, linear behavior at higher temperatures,

relative flatness at low temperatures, and zero

slope at T = 0 K. Both materials show slight

anomalies in the temperature dependence of

the bulk modulus. While Inconel X-750 shows

a concavity upwards in the B{T) curve, Inconel

718, as shown in Fig. 4, shows a maximum

near 100 K in its B(T) curve. Based on the

previous experience of this laboratory, it is believed in this case that the imprecisions in both Ci and Ct are a few parts in 10^. Thus,

this anomaly (though small) is beUeved to be

real and not a measurement-computational

artifact.

The interpretation of the maximum in the

B{T) curve of Inconel 718 can only be speculative. Unlike Inconel X-750, it was verified

during the present study that Inconel 718

is non-magnetic between room temperature

and liquid-nitrogen temperature (77 K). In some respects, the low-temperature behavior

of the bulk modulus of Inconel 718 resembles

that of stainless steels AISI 304, 310 and 316

reported previously [7]. In those cases, the

anomalous behavior was interpreted in terms of

the Doring effect that results from a large

volume magnetostriction accompanying a paramagnetic-to-antiferromagnetic transition.

However, no evidence of a low-temperature

magnetic transition in Inconel 718 seems to

exist; and neither the longitudinal modulus nor Young's modulus shows any evidence of a

dilatational anomaly. Thus, the nature of this anomaly, its magnitude and perhaps even the

question of its existence deserve further study.

Comparisons with some previous results [8]

show reasonable agreement for both the roomtemperature elastic constants and the temperature derivatives of E and G . It is easy to verify

that the temperature derivative of v reported

here is more reasonable than that obtained

from data in ref. 8. Differentiation of the

standard relationship

gives

1 dp ^ E (IdE ldG\

u dT E-2G\EdT GdT/

"

The present data satisfy eqn. (8) exactly, while

the data from ref. 8 give a L.H.S./R.H.S. ratio

of 2.6. Temperature derivatives of B have ap- parently not been reported previously. The

only previously reported elastic constant at 4 K is Young's modulus [2] ; this value, 2.11 X 10-^^N/m^, is identical with the present value.

This work was supported in part by the

Advanced Research Projects Agency of the

U.S. Department of Defense. Dr. E.R. Naimon

of Dow Chemical (Rocky Flats Division)

commented on the manuscript.

REFERENCES

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

Mater. Sci. Eng., to be published.

68
: Handbook on Materials for Superconducting

Machinery-. Battelle Metals and Ceramics Information Center, Columbus, Ohio. 1974.

3 F.H. Herbstein, Adv. Phys., 10 (1961) 313.

4 H.J. McSkimin, J. Acoust. Soc. Am., 33 (19611 12.

5 H.M. Ledbetter and R.P. Reed, J. Phys. Chem. Ref. Data, 2 (1974) 531.

6 Y.P. Varshni, Phys. Rev., B2 (1970) 3952.

7 H.M. Ledbetter. W.F. Weston and E.R. Naimon, J. Appl. Phys.. to be published.

8 Huntington Alloys: Inconel Alloy 713. Intern. Nickel Co., Huntington, West Virginia.

69
JOURNAL OF MATERIALS SCIENCE 10 (1975) 1309-1316 Low-temperature elastic properties of four

wrought and annealed aluminium alloys

E. R. NAIMON*, H. M. LEDBETTER, W. F. WESTON

Cryogenics Division, Institute for Basic Standards, National Bureau of Standards, Boulder,

Colorado, USA

The elastic properties of four annealed polycrystalline commercial aluminium alloys were

studied between 4 and 300 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), Poisson's ratio, and elastic Debye temperature.

The elastic stiffnesses of the alloys increase 4 to 13% on cooling from room temperature to

liquid helium temperature. The elastic constant-temperature curves exhibit regular

behaviour.

1. Introduction

Aluminium alloys are used extensively at cryogenic temperatures because of their favourable

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 for understanding its mechanical

behaviour. Most mechanical behaviour is best described by a dislocation model, and the elastic constants (usually the shear modulus and

Poisson's ratio) occur in most equations des- cribing the stress-strain state of a dislocated

solid. In this paper, the elastic properties of four

wrought aluminium 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 (4 to 13%), 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. An ultrasonic (10 MHz) pulse-superposition

method was used for determining the velocity, v, of a sound pulse propagated through the speci- men. 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 varia- tions of the elastic constants are typically a few

percent or less; and tests are completely nondestructive. Low-temperature elastic data for aluminium

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 behaviour of the alloys to be accurately

predicted between room temperature and liquid

helium temperature. Accurate elastic data

Present address: Dow Chemical USA, Rocky Flats Division, Golden, Colorado, 80401 © 1975 Chapman and Hall Ltd. 1 309

70
K. R. NAIMON. H, M. LEnBTiTTER, W. F WESTON

become especially important in applications

involving high stresses, large structural parts, or

precision parts.

2. Materials

Aluminium alloy 1100 is commercial quality

aluminium. 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. Aluminium 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 material. Aluminium alloy 7005 is a heat-treatable alloy

containing zinc and magnesium, which are balanced to obtain a natural-ageing alloy.

Chromium is added to reduce corrosion of the

heat-afTected weld zones, and zirconium is added to reduce weld cracking and to improve

mechanical properties.

Aluminium 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.

(However, the studies reported here were made

on annealed alloys.) Details of compositions, heat-treatments, mass densities, and hardnesses of the alloys are given in Tables 1 and II.

3. Experimental

Alloys were obtained from commercial sources;

1 100 and 7075 in the form of 1.9 cm rods, 5083

and 7005 in the form of 1.9 cm thick plate.

Cylindrical specimens 1.6 cm diameter and 1.6 cm long were prepared by grinding. Opposite

faces were flat and parallel within 2.5 |am.

Specimens were annealed at a pressure of

5 X 10 " Torr, or less, and cooled in the furnace. Hardnesses were determined by stan- dard 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 roomtemperature measurements and with a stopcock

grease for lower temperatures. In a few cases,

failure of these bonds at very low temperatures

required using a silicone fluid (viscosity = 2 x

I0-' P at 25 C) for bonding. The low-temperature

apparatus was described previously [I]. A pulse-superposition method was used to determine the sound-wave velocities over the temperature range 300 to 4 K. Details concerning

this method were given elsewhere [2].

4. Results

Quantities that were measured directly are the longitudinal and the transverse sound-wave

TABLE 1 Compositions of the alloys, mil 1 analyses, wl "„ Alkn Al C r Cu Fe Mg Mn Ni Si Ti V Zn Zr

1100 Bal 0.2 0.6 0.1 508.1 Bal 0.13 0.04 o.jy 4.75 0.63 0.003 0.08 0.01 0.007 0.04 7005 Bal 0.25 <0.l <0.4 1.2 <0.2 <0.3 <0.l 4.6 0.3 7075 Bal 0.3 1.6 0.7 2.5 0.3 0.5 0.2 5.6 TABLE II Properties of the alloys

Alloy Hardness (DPH no.. Mass density at Condition

1 kg load) 294 K (g cm ')

1 100 28 2.818 Annealed 345 C; furnace cooled

5083 78 2.666 Annealed 413 C, 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

1310

71
LOW-TEMPERATURE ELASTIC PROPERTIES OF FOUR ALUMINIUM ALLOYS

Figure I Temperature variation of the longitudinal modulus.

150 200 TCMPERATURE (K)

velocities v\ and ut- From these, the longitudinal

modulus, Ci, and the transverse modulus, G,

were calculated according to

and

Ci = pvf

Ct = pvt^

(1)

(2) These moduli are shown in Figs. 1 and 2 for the temperature range studied. No corrections were made for the change of mass density with

temperature; for aluminium this introduces a maximum error, over 300 K, of 0.4" Errors in the absolute velocities are believed to be about

1% or less. All 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 fl, and Poisson's

ratio V - are given by :

6 = C,

,

(3)

0.324

0.252 -

Figure 2 Temperature varialion of the transverse or shear modulus.

150 200 TEMPERATURE (K)

1311

72
E. R. NAIMON, H. M. LEDBETTER, W. F. WESTON

E = 3G(Ci - f Ct)/(Ci - Ct) , (4) 5 = Ci - f Ct , (5) and

V = i(Ci - 2Ct)/(Ci - Ct) . (6) The elastic constants obtained from these

relationships are shown as functions of temperature in Figs. 3 to 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 Ci and Ct were fitted to a

theoretical relationship suggested by Varshni [3]

:

C=C°- ' (7)

where C is any elastic constant (Ci and Ct 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 — sjtis the high-temperature limit of the temperature derivative dC/dr. 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.

Figure 3 Temperature variation of Young's

modulus.

!50 ?no ^tMPFRATURC (I.)

0 50 ; Hi ISO ?no 250 "no TEMPEHATURF (K) 1312

73
LOW-TEMPERATURE ELASTIC PROPERTIES OF FOUR ALUMINIUM ALLOYS

o

_3>

rrj — <N O m m m d d d

—. rr

o\ o fN rJ m

— 0\ <N

r- OS

r~- r-^ r~ d d d

lO —

>

r*^ vo vo

00 p O ON O fN

d d d d d d d d d d d d d d d d d d d d

osrnioO\oovo — r-(Noo — ^u^ovtNvoov^oom inr-~oooovooooo^t^o\os^r^O\osy3r— ooo\

d d d d d d d d d d d d d d d d d d d d

-TOOooOS'OOO^Ov"/^'^ 0\r-<Nr<-ir<-i — OO-^rr

d d d d d d d d d d d d d d d d d d d d

T}-ooo\ — v£)oso\(^ia\'/^r-o r^oooo — — r-o-* O — <Nm»or^o\Ov

OOoOOOOOOOOOOOOOOOOO

S§°888^S88^8S8^SSS° fN _ r«-i (N^ m(N-^ rnfS^ m(N«

8 8

B • 2 6 < .2

a

— <u

03 T3

00 u

HI

=23 .a U W

1313

74
E. R. NAIMON, H. M. LEDBETTER, W. F. WESTON

Parameters C\ s, and t (determined by a least- squares fit of Equation 7 to the data) are given

in Table IV. Average differences between

where

TABLE 1 V Parameters n Equation 7

Alloy Mode C" (10" N m-) i(10"Nm-^) t (K)

1 100 1 1.136 0.0903 255.9

t 0.2892 0.0301 206.9 5083 ! 1.151 0.0926 235.3

t 0.3067 0.0381 206.7 7005 1 1.175 0.0928 240.0

t 0.2963 0.0327 203.4 7075 1 1.126 0.0707 219.6

t 0.2961 0.0344 223.3

measured and curve values are 0.03% and 0.06%

for the longitudinal and transverse moduli,

respectively. 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.

TABLE V Temperature coefficients of the elastic constants at room temperature (10" " K

1 dfl 1 d£ 1 dC 1 d r Alloy B dT E dT G dT dr

1 100 -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 Aluminium* — 1 .97 -5.53 -5.77 1.80 *Derived from single-crystal data in [4].

The elastic Debye temperature, 9, can be

calculated from the elastic wave velocities by [5] :

(8)

K

k \4^/ (9)

Here h is Planck's constant, k is Boltzmann's

constant, is Avogadro's constant, p is the mass

density, and A is the effective atomic weight. The

average velocity is given by + 2vt- -1/3 (10)

The elastic Debye temperatures for the four

alloys at r = 0 K, and also for unalloyed

aluminium are given in Table VI.

TABLE VI Elastic Debye temperatures at T = 0 K

Alloy 9(K)

1100 426.2 5083 440.4 7005 425.9 7075 422.2 Aluminium 430.6*

*Calculated from single-crystal data in [4].

5. Discussion

As shown by the data in Table III, changes in the

elastic constants of aluminium and its alloys

between 300 and 4 K are about 4% for B and

V, and 12% for E and C. 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 temperature. Thus, any

changes in the mechanical behaviour of these 1314

75
LOW-TEMPERATURE ELASTIC PROPERTIES OF FOUR ALUMINIUM ALLOYS

alloys in this temperature region probably cannot

be ascribed to an elastic origin.

The temperature behaviour of the elastic constants of the aluminium alloys reported on

here is quite regular. The moduli decrease

regularly with increasing temperature and the modulus-temperature curves are relatively flat at low temperatures. Also, in accord with the third law of thermodynamics, the slopes dC/dT

approach zero at zero temperature. A linear temperature dependence is exhibited above about

150 K, which is roughly one-third of the Debye

temperature. Besides indicating the absence of

magnetic or structural transitions, this ideal temperature behaviour suggests that the an- harrnonic 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. Aluminium single crystals are only slightly anisotropic; the Zener anisotropy

ratio for aluminium is 1.2; it is 1.0 for the

isotropic case. Aluminium's isotropy is purely

accidental since aluminium, because of its three valence electrons, has a large band-structure

contribution to the elastic constants [10]. This

contribution is usually anisotropic; the anisotropy is cancelled in the case of aluminium by

other energy terms. Thus, since aluminium single

crystals are only slightly anisotropic, even a

strongly textured polycrystalline aggregate of

aluminium would have nearly isotropic elastic behaviour.

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 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 formulae given by Landau

and Lifshitz [6] it can be shown foraluminium at room temperature:

0.005

Bt = 0.045 (11)

= 0.020 , and — 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 \.

Effects of alloying on the elastic properties of

aluminium cannot be accurately determined

from the present study because of the large

number of alloying elements and their inter- actions. Such effects have been considered

elsewhere [11]. However, some general observations can be made concerning alloying. Disregarding alloy 1100, with respect to '.'pure" aluminium, 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

1 100 alloy, which contains only 1°,, of impurities,

presents an interesting case. The data indicate th^t while the shear modulus of this alloy is identical to that of unalloyed aluminium, the bulk

modulus is higher by about 5%. Since the bulk

modulus is not measured directly, but is cal- culated 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 the alloys

and for all temperatures, as a first approximation.

and

B ^ 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.

6. Conclusions

From the results of this study the following

conclusions are drawn:

(1) all the elastic properties of aluminium

alloys 1 100, 5083, 7005, and 7075 behave regularly

with respect to temperature;

(2) for all alloys studied, the temperature

behaviour of both Ci and C'l can be described

accurately by a theoretical relationship suggested

by Varshni;

(3) in this series of alloys, in the annealed

condition, 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.

1315

76
E. R. NAIMON, H. M. LEDBETTER, W. F. WESTON

Acknowiedgement

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 (1974) 246.

2. H. J. MCSKIMIN, /. Accoust. Soc. Amer. 33 (1961)

12.

3. Y. p. VARSHNi, Phys. Rev. B2 (1970) 3952.

4. G. N. KAMM and o. A. ALERS, /. 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" (ASM,

Metals Park, Ohio, 1961).

8. R. DEVELAY, A. FAURE, S. LEHONGRE, D. MUGNiER and D. SCHROETER, in "Advances in Cryogenic Engineering", Vol. 12, edited by K. Timmerhaus (Plenum, New York, 1967) p. 484.

9. J. L. CHRISTIAN and J. F. WATSON, ibid. Vol. 6, (1961) p. 604.

10. T. SUZUKI, Phys. Rev. B3 (1971) 4007.

11. w. KOSTER, Z. Metallk. 32 (1940) 282.

Received 11 December 1974 and accepted 28 January

1975.

1316

77
Low-temperature elastic constants of a superconducting coil

composite

W. F. Weston*

Cryogenics Division, Institute for Basic Standards, National Bureau of Standards, Boulder, Colorado 80302

(Received 27 Feoruary 1975; in final form 17 June 1975) A resonant piezoelectric oscillator method for measuring elastic moduli was applied to composite materials. The complete set of elastic compliances of a superconducting coil composite was determined

semicontinuously between 4 and 300 K. Also, two moduli of a layered fiber-glass-epoxy composite were

determined; this composite is essentially the matrix material of the coil composite. The Young's moduli,

shear moduli, Poisson ratios, and elastic stiffness coefficients are also reported. Results agree closely with

elastic data obtained by conventional testing methods. PACS numbers: 62.20.D, 74.50.

I. INTRODUCTION

The elastic properties of composite materials are

currently of considerable interest. Many composites,

because they have high mechanical strength and a high

modulus-to-density ratio, are useful as strong lightweight structural materials. Some composites have

been designed for use at low temperatures, for example, as superconducting coils for high-field magnets.

However, few elastic data are available for composite

materials, particularly at low temperatures. To the author's knowledge, no complete sets of elastic data are available for composite materials at 4 K. Some data are useful as design parameters, and elastic constants are among the most accurately known fundamental

physical properties of solids. Because composite materials are usually highly

anisotropic, the determination of their elastic constants

is much more difficult than for the more usual quasiisotropic engineering materials. Conventional methods

of measuring the elastic properties of solids have been

applied to composites with only limited success. These

methods usually require relatively large specimens and considerable time and effort, especially if a complete

set of elastic constants is required.' These methods, as applied to composites, include vibrating^'^ or resonant- beam tests, tensile or compressive tests, '•''•^ torsion

tests, ^'^ bending tests, '•'° and pressure tests. ^'"'^ Generally, the elastic constants measured by these methods

are less accurate than those acquires by ultrasonic

tests.

Ultrasonic pulse techniques are limited by the necessity of using wavelengths that are large compared with

fiber diameters and at the same time sufficiently small

compared with the dimensions of the specimen so that true plane-wave conditions can exist. Since ultrasonic

pulse methods are typically performed at megahertz

frequencies, this condition cannot be met for many

composites. If the wavelength is not larger than the

fiber diameter, the wave is attenuated and scattered. Ultrasonic pulse-echo methods have been applied to some metal-matrix composites with extremely small

fibers.""'" Even in these cases the pulse-echo trains are poor because of high attenuation and dispersion.''

In fact, Achenbach and Herrman" predicted large dispersive effects for shear waves propagated in the fiber direction, even when the wavelength is much larger

4458 Journal of Applied Physics, Vol. 46, No. 10, October 1975

than fiber diameters. Zimmer and Cost" verified this prediction by measuring sound velocities in a unidirectional glass -reinforced epoxy-fiber composite. For 1 wavelength-to-fiber diameter ratios of roughly 50, dispersion increased some phase velocities by a factor of !' 2. Elastic constants of some carbon-fiber-reinforced

plastics have also been measured by ultrasonic pulseecho techniques. '^~^'

Attempts to use 10-MHz pulse-echo techniques on the composite studied herein were unsuccessful. Both longi- < tundal and shear waves were propagated along all orientations and produced no detectable echoes. This is understandable since the filament sizes for this composite, i, as described in the text, are larger than the sound wave- ; length. Larger specimens than are currently available . would be needed for 1—5-MHz pulses, although it is doubtful that even these frequencies would produce useable echo patterns for this material.

The ultrasonic immersion technique, as described by

Markham," is fast and accurate and offers the possibility of measuring all the elastic properties on one specimen. However, this technique is also limited by

fiber size and by sample dimensions.^^-^* The choice of ultrasonic modes that can be propagated in a given direction in a specimen is also limited, and the technique is inapplicable to low-temperature measurements.

Zecca and Hay^' avoided the problems encountered

with megahertz frequencies by using electrostatic

^ transducers to generate and to detect resonant kilohertz frequencies of a metal-matrix composite. Electrostatic

transducers, however, cannot be used with nonmetallic ^ materials. ''

6 The most widely used ultrasonic technique, using i kilohertz frequencies, is the resonant piezoelectric > oscillator. This method is due principally to Quimby

and associates. Quimby thoroughly analyzed the vibra- j tions of solid rods driven by a piezoelectric crystal, and | Balamuth^'' first employed drivers and specimens of matching fundamental frequencies. Rose^* extended this method to torsional oscillators. Basically, the piezoelectric properties of quartz are used to generate and to detect resonant frequencies of an oscillator consisting

of one or two quartz crystals, a specimen, and perhaps

a dummy rod, all cemented together. This method has

been used to measure elastic moduli at both high^^-^" and

Copyright © 1975 American Institute of Physics 4458

78
RESONANT PIEZOELECTRIC OSCIlLtlOR

FIG. 1. Resonant piezoelectric oscillator apparatus.

low^''-^^'^°'^^ temperatures, and has worked well with both metals^^'" and nonmetals.^''"^^ Thus, this technique is well understood both theoretically and practically, and it would seem to be well suited for testing composites.

The long wavelengths used (t\'pically on the order of centimeters in the kilohertz region) are much larger

than fiber di:imeters, and, since the technique employs

standing waves rather than traveling waves, specimen

sizes can still be kept small.

In the p^e^e^t paper it is described how this resonance

technique was used to measure the complete set of elastic compliances of a superconducting coil composite,

which was essentially composed of an epoxy matrix and

unidirectional copper-niobium-titanium fibers. The

compliances were measured semicontinuously from 300

to 4 K. Results are given also for an epoxy—fiber-glass

layered composite. The results generally agreed closely

with the few existing static measurements on the same

composite.

II. EXPERIMENTAL PROCEDURE

A. Resonant piezoelectric oscillator The three-component resonant piezoelectric oscillator technique was originally described by Marx,^* and was

discussed in detail by Fine.^^ The technique consists of bonding quartz-driver and quartz-gauge piezoelectric

crystals to a specimen to produce and to detect a

standing longitudinal (or torsional) resonant wave. Each

component's length is adjusted so that its resonant fre- quency is closely matched to that of the other components. The system is then driven by the driver trans- ducer at its resonant frequency, which is monitored by

the gauge crystal.

Longitudinal waves are excited by a -quartz bars of square cross section, excited into longitudinal vibration by an ac signal applied to full-length adherent electrodes (see Fig. 1). The driver quartz was fully gold

plated on two parallel sides {z faces') with a shallow

notch cut in the center of each electrode. Fine gold445? J. Appl. Phys.. Vol. 46 Mo. 10, October 1975

plated tungsten wires, seated in these notches, served i to suspend the resonator assembly at displacement j nodes and provided the necessary electrical contact. ; The gauge quartz was gold plated only over the center

third of the two parallel faces and electrical contact was made through fine wires adhered into the notches j with silver conducting paint. The electric axes of the driver and the gauge crystals were perpendicular to reduce electrical pickup. j The longitudinal oscillator apparatus is illustrated i in Fig. 1. The driver-quartz crystal was suspended by ; 0.003 in gold-plated tungsten wires. The driver mounts,

which were aluminum, were removed during the tor- ! sional measurements, since the torsional driver crystal I was suspended between the Bakelite posts by nylon

j threads. The specimens were cylindrical for all mea- i surements. The chromel-constantan thermocouple was

j placed near, but not touching, the specimen. The en- , tire apparatus was sealed in a thin-walled stainless- ' steel can. The can was partially evacuated, placed in the ullage of a helium Dewar, and lowered stepwise to achieve cooling. ^ The torsional quartz crystals were circular with J their length in the .r direction and with four gold-plated ^ electrodes running the length of the crystal. Opposite I pairs of electrodes were electrically connected. The j driver quartz was suspended vertically by means of fine

j nylon treads attached with varnish to opposite sides of | the quartz near a displacement node of vibration. J

i The quartz crystals were bonded together with a semipermanent adhesive, Eastman 910.^' The specimens were right-circular cylinders and were also bonded to the gauge quartz with the same adhesive. For some i torsional experiments, this cement occasionally failed

j at low temperatures; such experiments were repeated .! until successful. Other materials such as vacuum | grease and epoxy resin have been used as low-tem- | perature bonding agents

.

The quartz crystals used had resonant frequencies of ! 60 and 100 kHz. Most of the measurements were done ; with the 60-kHz quartz crystals (^-in. square cross

j for longitudinal, and ^ in. in diameter for torsional),

j but some were done with the 100-kHz crystals {| in. in diameter). The mass of each driver-gauge combination ] was noted and the resonant frequencies were monitored

from 300 to 4 K. The length of each specimen was

determined such that its resonant frequency for the entire temperature range was within approximately 5% of the oscillator assembly although in some cases the specimen frequency differed by as much as 10%. The

mass, length, and diameter of each specimen were

noted and its mass density was determined by Archi- | medes's method using distilled water as a standard.

The specimen was then cemented to the quartz crystals and the resonant frequency of the oscillator was monitored from 300 to 4 K. No thermal contraction corrections were made; for the coil composite described be- ' low this introduces a maximum error (over a 300 K ] range) of about 0.5%.^^ Maximum uncertainties in the frequency measurements are estimated to be about one ' part in 10^.

W.F.Weston 1459 j

79
8. Electronics TABLE I. Dimensions and orientations of specimens.

Most workers who have used the piezoelectric oscilla- tor technique have employed electronics similar to those

described by Marx.^** This system requires manual ad- justment of the input frequency to the driver crystal to keep the oscillator at resonance. A less time-consuming

and more accurate system is the regenerative system

designed by Gerk"*" and slightly modified by Johnson.

Gerk's system consists primarily of a closed loop

containing an amplifier and the driver-gauge combination of piezoelectric crystals, which act as a resonant element, with a feedback circuit to control the gain of the amplifier. This system is based on the principles

of an oscillator; that is, the output of an amplifier is fed into the input through a frequency-selective network

(the quartz crystals). The phase shifts are corrected

with a variable phase shifter. A more complete description of the system was given by Gerk."*" This sys- tem locks onto the resonant frequency and automatically

stays at the resonant frequency if this frequency

changes, due to a change of the specimen's environment, for example.

C. Specimens

The superconducting coil composite is fully described elsewhere.^® Basically, the coil is composed

of a copper-stabilized niobium -titanium wire (0.56

xO.72 mm cross section) coil impregnated with epoxy. The wire layers are separated by layers of dry fiber- glass cloth (0.1 mm thick). The copper-to-superconductor volume ratio of the wire is 1.8-1. The wire-toepoxy cross-sectional area fraction is about 3 to 1. The

coil dimensions were approximately 19 cm i.d. , 24 cm o.d. , and 10 cm long.

Coordinate axes were chosen to coincide with the specimen axes. The longitudinal axis (3 axis) coincides

with the axis of the wire, which was assumed to have

zero helix angle. The 1 and 2 axes lie along the radius

and coil axes, respectively, and form an orthogonal

set with the 3 axis. The symmetry of the coil (orthorhombic) requires nine elastic constants to characterize

the material. By assuming the material to be trans- versely isotropic and by neglecting the effects of curvature, the number of independent elastic constants is re- duced to five. An orthotropic body with transverse iso- tropy has the same symmetry as close-packed hexagonal crystals such as magnesium and zinc. Justification for equating the 1 and 2 axes is given elsewhere,

and additional experimental data show that the effect of curvature on the moduli is minor.

The masses, lengths, and angles between the speci- men axis and the 3 axis are given in Table I. The density of the coil composite was found to be 6.000 g/cm^.

Besides the superconducting coil, some specimens

composed only of the fiber-glass cloth and epoxy were

available. The fiber-glass—to— epoxy volume ratio in these samples was approximately the same as in the

cuil composite. Measurements were made on two speci- mens of this material. Dimensions and orientations of these specimens are also reported in Table I, with the angle 9 corresponding to the angle between the speci4460 J. Appl. Phys., Vol. 46, No. 10, October 1975

Specimen Diameter Mass Length 6 No. (in.) (g) . (cm) (deg)

|| Superconducting coil composite specimens !|

1 0.1250 0. 540 1,1128 90 j 2 0.1250 0.925 1.9439 0

3 0.1863 1.081 1.0818 0 4 0. 1890 1.248 1.1481 90

5 0.1875 1. 874 1.7600 60 I 6 0.1851 1,327 1,2555 60 Epoxy-fiber-glass composite specimens

1 0, 1875 0. 885 2.5387 0

2 0. 1878 0, 520 1,4862 0

men axis and the direction parallel to the layers of fiber-glass cloth. The density of this material was

found to be 1.931 g/cm^. This material can also be

considered transversely isotropic, if the fiber-glass

is considered transversely isotropic. Thus, five constants are again needed to completely characterize the material elastically. Cylindrical specimens of only one

orientation were available; thus, only two elastic constants of the fiber-glass-epoxy matrix material were

determined. '

(II. RESULTS

The directly measured quantity in these measurements is /o, the frequency of the three-component piezo- electric resonator. For longitudinal resonance the fre- quency of the specimen (f^) can be found from f^, the frequency of the driver-gauge quartz assembly {fj, the mass of the specimen {mj, and the mass of the quartz

(m,) «:

fl=fl + in-fM,/m^. (1) This formula is more exact than the approximation

/.=/o + (/o-/>/w,, (2) which is usually quoted and used for resonant

oscillators.^'* The formula for the torsional oscillator is somewhat ' more complicated since the moments of inertia (rather

than the masses) of the components are involved^^: fs=fl + ifo-fl)m,7i/my,. (3) Here, is the radius of the quartz and is the radius

of the specimen.

The Young's modulus E or shear modulus G for the particular orientation of the specimen is given by

E or G = 4L^flp. (4) E is found from the longitudinal-mode fundamental fre- quency and G is found from the torsional -mode fundamental frequency. In Eq. (4), L is the specimen length

and p is the mass density.

Thus, the measured quantities are the Young's and shear moduli, which are directly related to the elastic compliances S^J. However, if the wavelength is not much larger than the sample dimensions, a correction

for the Poisson contraction (or lateral motion) must be

W.F. Weston 4460

80

Ii
composite.

Exp.

No.

Spec. No. Mode e Measured quantity

1 1 L 90

2 2 L 0

^33

3 3 T 0 ^13= T~

4 4 T 90 c

^

^44 + '^66

5 5 L 60 r

9^11 '^33 6^13 3S44

6 6 T 60

8 SSjj + 3S33 — 65i3 + 2S44 + SSee

^plied to Young's modulus. When the length-to-diameter ratio is 10 to 1 , this correction amounts to about one part in 10^. The value of Young's modulus for an arbitrary direction in a specimen with transverse isotropy is given in terms of the elastic compliances by^

1 /£ = s in* 61 + S33 cos^e + (2S13 + S44) sin^ e cos^ 6 , (5

)

where 6 is the angle between the specimen axis and the unique axis. Similarly, the shear modulus is 1/G = S« + (Si, - - iSii) sin^e + 2 (Sii + S33 - 2Si3 - S44) sin^ 6 cos^ 6

.

Most measurements were made during cooling. All experiments were performed at least twice, and any unusual behavior was examined on heating as well as

cooling, for reversibility.

A. Superconducting coil composite

Measurements made on the coil composite are listed

in Table n. Again, 6 is the angle between the specimen

axis and the 3 axis (longitudinal axis). The mode specifies either longitudinal or torsional oscillations. The

7.972 I 1 1 1 1 1 1 1 0 50 100 150 200 250 300 350

TEMPERATURE (K) FIG. 3. Young's modulus of the superconducting coll composite, e = o\

subscripts on E and G designate the direction of the force and the plane on which the force is ^plied, respectively. The compliances S^, S33, and are found

directly from experiments 1, 2, and 3, respectively.

From experiment 4, S^^ found using the relation 2(Sn -Si2)==Sge and the value of S^^ from experiment 3. Either specimen 5 or 6 can be used to determine S^^. This compliance must be determined from a measurement on a specimen with its axis at an angle to the

longitudinal direction. As can be seen from Table O,

however, the Young's or shear modulus measured on specimens with 0° < 6< 90° is only partly dependent on

. Thus, Sj3 will be the least accurately determined

compliance. The measurement of this compliance cor- responds to the measurement of the elastic stiffness by megahertz ultrasonic techniques. Zimmer and Cost" encountered difficulty in measuring C^^ for a composite material and their estimated uncertainty in this elastic constant (100%) was ascribed to dispersion

as well as to the relatively high inaccuracy.

The experimental data are presented in Figs. 2—7.

The smooth curves represent regular temperature beso 100 150 200

TEMPERATURE (K)

250 300

FIG. 2. Young's modulus of the superconducting coll composite, 6 = 90°.

100 150 200

TEMPERATURE (K)

250 30O

FIG. 4. Shear modulus of the superconducting cotl composite,

6 = 0°.

4461 J. Appl. Phys., Vol. 46, No. 10, October 1975 W.F. Weston 4461

81
150 200

TEMPERATURE (K)

FIG. 5. Shear modulus of the superconducting coil composite,

6 = 90°.

TEMPERATURE (K)

FIG. 7. Shear modulus of the superconducting coil composite,

6 = 60°.

havior as given by a theoretical relationship suggested

by Varshni*^ that has been used successfully to represent the temperature dependences of metals,*^ ionic

solids,'*^ and several alloys.''® The over-all fit of this function to these data is not nearly as good as for other> materials, and the data for £33 could not be fitted to the Varshni function because of its concave-upward curvature. A smooth curve was drawn through the data points

and the actual point values were used in calculations. This modulus was measured several times to check its behavior, and data from different measurements agreed

closely. These data had completely reversible temperature response

The shear modulus corresponding to shearing the

axial planes in the radial direction G^j, or vice versa,

is shown in Fig. 8. This modulus was calculated from

the curve values of G13 and G', since Gi2 = l/S66- Because of the difficulty in machining the specimens

to small radii, the ratios of lengths to diameters were

in some cases on the order of 4 or 5 to 1. The estimated

correction to the Young's moduli in these cases is still only about 1%.

150 200

TEMPERATURE (K)

The elastic compliances as functions of temperature

are given in Table HI. The compliances S^, S33, and

S44 are just reciprocals of S^, £33, and G13, as mentioned previously. The error in the directly measured

moduli is estimated to be about 1%; thus these compliances should be accurate to 1% also. The complianci

is found from combining the results of three measurements. Thus, may be less accurately known. It was found that in solving for S^^ from either Eq. (5) or '

(6), a 1% error in the measured modulus led to a 10—

20% error in S13. Thus, the error associated with 5^3

is considerable, perhaps as much as 100%. The values reported for S^^ are the average of the values found * from experiments 5 and 6, which are -0.110 and

-0.058, respectively. From the material available, it 1 was not possible to machine specimens with 6< 60° for other determinations of Sjg. As expected, the material was found to be highly anisotropic with Sn = 2s33.

It is emphasized that the data reported here are dynamic (adiabatic) rather than static (isothermal) and

apply to rapid, rather than slow, loading. Conversion

formulas are given in Landau and Litschitz,*' for ex150 200

TEMPERATURE (K)

FIG.

Site,

6. Young's modulus of the superconducting coil compo6 = 60'. FIG. 8, Shear modulus Gjj of the superconducting coil composite.

4462 J. Appl. Phys., Vol. 46, No. 10, October 1975 W.F. Weston 446:

82
TABLE III. Elastic compliances of the superconducting coLl composite at selected temperatures (10'"^ m^/N) as determined

bv resonance oscillator tecliniques. T (K) ll c

12 e

13 c o c

0 0. 242 -0,074 -0.061 0. 112 0,910 0.631

50 0,242 -0.074 -0,061 0. 113 0,910 0, 632

100 0,248 -0.076 -0,066 0. 115 0,920 0,648

150 0. 260 -0,083 -0.073 0.118 0. 964 0, 684

200 0,275 -0. 0S9 -0.079 0.119 1.046 0,728

250 0.294 -0.094 -0,082 0.121 1. 170 0, 776

300 0.316 -0.096 -0.084 0,123 1.350 0, 823

ample; in most cases the differences between adiabatic and isothermal elastic constants are small. These

formulas involve the thermal-expansion coefficients and

the specific heat. The thermal -expansion coefficients are known for this composite, but the specific heat has

not yet been determined. If an estimate of the specific heats of the components, the difference between the

adiabatic and isothermal Young's moduli is at most

0.5%. The adiabatic and isothermal shear moduli are,

of course, equal,

B. Epoxy-fiber-glass composite

Measurements were also made on the epoxy-fiberglass specimens listed in Table I. These specimens were oriented with the layers of fiber-glass cloth running the length of the specimen; these were the only

orientations available. The Young's and shear moduli

are shown in Figs. 9 and 10. The shear modulus is well represented by the Varshni function, and is seen to have

normal temperature behavior. A smooth curve was

drawn through the Young's modulus data, which could

not be fitted to the Varshni function because of the lowtemperature maximum.

IV. DISCUSSION

Some elastic data, taken by conventional static methods, were available for the superconducting coil composite and the epoxy-fiber-glass composite.

These data generally compared very favorably with the resonance data. For example, resonance values of E^^

at room temperature and 4 K are 12 and 1% larger than

3.096 -

0 50 100 150 200 250 300 350

TEMPERATURE (K)

FIG, 9, Young's modulus of epoxy—fiber-glass, 6 = 0°.

4463 J. Appl. Phys., Vol. 46, No. 10, October 1975

83

the static data taken at these temperatures. The inaccuracy of the static values is given as about 10%. Also, resonance values of £33 agree with the static data within 5—7%, and resonance values of G^^ agree with the static values within 6—11%,

Other elastic constants of interest include the Poisson

ratios: Vn'-^-SjS,, (7) and

"^is/Saa, (8) where v ^j represents the negative ratio of strain in the

j direction to strain in the i direction. The value of room temperature is 0.304. The

reported^^ static value is 0.335, with the inaccuracy

again believed to be about 10%. Thus, this result gives

added confidence to the resonance value of Sjg. However,

the Poisson ratio v^^, involves the compliance S^^, which has already been noted to be much more inac- curate than the other four compliances. In fact, if the room -temperature values of and S33 are used, 1/^3 = 0.683, whereas the static value is 0.333.^^ Since the other elastic constants are all in relatively good agree- ment with the static values, it is reasonable to expect

that j/j3 should be also. Thus, the value of 5^3 is highly

suspect, especially since a value near j seems more

realistic for Poisson's ratio.

It is also of interest to consider the compressibilities

.

For a material with transverse isotropy, the linear compressibilities are

fei = Sn+Sx2+S,3

and (9)

where the subscripts designate the axis. Also, the vol- ume compressibility is K = 2k^+k^. (10)

If the room-temperature values of the compliances,

as given in Table HI, are used, then fei = 0. 136x10"^°

0.940

I

1 1 \ 1

i r

0.904 -

0.61b -

0 50 100 150 200 250 300 350

TEMPERATURE (K)

FIG, 10, Shear modulus of epoxy—fiber-glass, 9 = 0°.

W.F.Weston 4463
TABLE IV. Elastic stiffnesses of the superconducting coil composite at selected temperatures (10^" N/m^). T (K) 11 C ^

12 C *13 C *

44 QGG

0 5. 10 1.94 2.39 10. 59 1. 10 1. 58

50 5.10 1.93 2.39 10.46 1. 10 1. 58 100 4. 99 1.90 2.34 10.27 1.09 1. 54 150 4. 80 1. 87 2. 27 10.05 1.04 1.46 200 4. 54 1. 79 2. 15 9. 86 0. 96 1. 37 250 4. 21 1.63 1.99 9.64 0. 85 1.29

300 3. 82 1.39 1.77 9, 35 0. 74 1.22

^These constants are dependent on 1^13, and were calculated using the static value of ^13, which was taken as independent

of temperature.

m2 N, /^3= -0.045x10-'° mVN, and A'= 0. 027 x 10-'° mVN. However, if the Poisson ratio 1^13 is taken as ^, then Si3 = - 0. 041 xlO"'° m^/N at room temperature, and

fei = 0.179x10-1° mVN feg = 0. 041 x lO-'" mVN, and the volume compressibility /s: = 0. 399 xlO"i° m^/N. The

relatively large value for and the negative ^3 that result from the resonance S13 may not be unreasonable,

considering the multicomponent material. The elastic stiffness coefficients, however, can be calculated from

the compliances and indicate that the resonance value of

S,o is in error. The elastic stiffnesses (C^j) can be directly calculated by inverting the S^j matrix ii all the compliances are known. Tiie explicit relationships are^"

1^-^.5 _s' )s

^13

c'-'44 = —O^44

-^11 + -big s

1

(11)

and

where § = 533(511 +Si2)-2S^3. Except for C^^ and Cgg, these relationships depend strongly on 5,3 . Using the resonance value of 5^3, the elastic stiffnesses that depend on 5i3 are found to increase with increasing temperature, which is contrary to normal behavior. Since

the Young's and shear moduli all behave fairly regularly

with temperature, the resonance value of 5^3 seems unreasonable. In calculating the elastic stiffnesses, then, a value of j for 1^,3 was used, and this Poisson's ratio was taken as constant with temperature.^® It is believed

that the elastic stiffnesses obtained this way are more

accurate than those obtained using the resonance value

of S13. The elastic stiffnesses are given in Table IV, As expected, the elastic properties of this composite

are highly anisotropic. The anisotropy is best described

by considering the percent elastic anisotropy as discussed by Chung and Buessem.'*® For transversely iso- tropic symmetry, this scheme involves using the conventional anisotropy ratios (the linear compressibility

ratio \ = k^/k^, and the shear anisotropy A^ = C^^/Cgg)

to calculate anisotropy parameters A* and A* with the following properties: A* is zero for materials that are

elastically isotropic, i. e. , A = 1; A* is always positive and a single-valued measure of the elastic anisotropy of a material regardless of whether A <1 or A > 1 ; and A*

gives the relative magnitude of the elastic anisotropy

present in the material. For this composite A^=0.23,

A* = 8.4%, A^==0.61, and A* =12.3%. For comparison,

A* « 9 . 9% and A* = 13 . 7% for zinc , '•^ which is one of the most anisotropic hexagonal metals.

The static value of Young's modulus of the fiberglass— epoxy is also available.*" This Young's modulus was measured on a specimen oriented like the one used

for the resonance measurements. The resonance and

static values differ by 4—10%. There was no evidence of a maximum in the static data, but it is believed that the maximum is real. This maximum was observed during

several measurements and was completely reversible.

No explanation for this behavior can be given at this time.

V. CONCLUSIONS

The resonant piezoelectric oscillator method used in this work is generally suitable for measuring the elastic properties of composite materials. This method is accurate, fast, and requires only small specimens. It is suitable for measurements at all temperatures from 0 K

to greater than 1000 'C.^^'^" The general techniques and theory of this method are well understood. This method

does not appear to have the limitations of static and

other ultrasonic techniques.

The elastic compliance data presented in this work

generally agree well with existing data taken by conventional static methods, with the exception of 5i3. It is believed, however, that if specimens with different orientations were available, a more accurate determination of could be made. These data can be used to cal- culate other elastic properties, such as the elastic stiff- ness coefficients and compressibilities. However, many

of these elastic properties depend on S13, and their ac- curacies would also benefit from a better determination

of S13. The superconducting coil composite was highly anisotropic, with elastic anisotropics similar to those of zinc. This high anisotropy, characteristic of many composite materials, necessitates careful measurements of elastic properties of these materials.

The measured elastic moduli of the composite are not as well represented by the Varshni function as they are

for most metals and alloys; i.e. , the elastic moduli of the superconducting coil composite are less regularly

behaved with respect to temperature.

ACKNOWLEDGMENTS

The author would like to express his graditude to D. L. Johnson for his considerable assistance with the piezoelectric oscillator techniques. He would also like to

4464 J. Aopl. Ph/ 46, No. 10, October 1975 W F. Weston 4464

84
in the experiments. The author is also indebted to H.M.

Ledbetter for many helpful discussions on elastic properties in general and for a critical reading of the manuscript. This work was supported in part by the Annapolis

l^aboratory of the Naval Ship Research and Development

Center and in part by the Advanced Research Projects

Agency of the U.S. Department of Defense.

'XRC-NBS Postdoctoral Research Associate, 1974—5.

P, R. Goggin, J. Mater. Sci. 8, 233 (1973).

-T.J. Dudek, J. Compos. Mater. 4, 232(1970).

T. HiraiandD.E. Kline, J. Compos. Mater. 7, 160(1973).

'Albert B. Schultz and Stephen W. Tsai, J. Compos. Mater.

2, 3(is (1968). 'R.C. Rossi, J.R. Cost, and K. R. Janowski, J. .Am. Ceram.

Soc. 55, 234 (1972). G.C. Wright, J. Sound Vib. 21, 205(1972).

R. E. .Allred and W. R. Hoover, J. Compos. Mater. 8, 15

(1974). ^L.J. Cohen and O. Ishai, J. Compos. Mater. 1. 1030(1967).

^N. L. Hancox, J. Mater. Sci. 7, 1030(1972). T.R. Tauchert, J. Compos. Mater. 5. 456(1971).

A. Smith, J. Phys. E 5, 274(1972).

-A. Smith, W. X. RevTiolds, and N. L. Hancox. J. Compos.

Mater. 7, 138 (1973). A. Smith, S.J. Wilkinson, and \V. .v. RevTiolds. J. .Mater.

•>ci. 9, 547 (1974).

•.V. Grabel and J. R. Cost, Met. Trans. 3, 1973(1972). A.E. Lord and D. Robert Hay, J, Compos. Mater. 6, 27-5 1972).

A'. Sachse, J. Compos. Mater. 8. 37-^(1974). .O. Achenbach and G. Hermann. ALAA J. 6. 1>)32 (196«).

;. E. Zimmer and J. R. Cost, J. .Acoust. Soc. .Am. 47, "95 (1970). A.N. Reynolds and S.J. Wilkinson, L'ltrasonics 12, 109

'1974).

S.J. Wilkinson and W.N. Reynolds, J. Phvs. D 7, 50

' 1974). -'T.R. Tauchert and .A. N. Guzelsu, Trans. .ASME 39, 9b '1972).

- Robert S. Smith, J, Appl. Phys. 43, 2555 (1972). "A.R. ZeccaandD.R. Hay, J. Compos. Mater. 4, 556

(1970).

"''S. L. Quimby, Phys. Rev. 25, 558(1925).

'L. Balamuth, Phys. Rev. 45, 715(1934). 27^^F.C. Rose, Phys. Rev. 49, 50 (1936). Hunter and S. Siegel, Phys. Rev. 61, 84 (1942).

J. W. Marx and J. M. Silvertsen, J. Appl. Phys. 24, 81 30^(1953). ^^M. E. Fine, Rev. Sci. Instrum. 25, 1188(1954).

49, 663 (1936).

J. Zacarias, Phys. Rev. 44, 116(1933).

J. Marx, Rev. Sci. Instrum. 22, 503 (1951).

^^S. Siegel and S. L. Quimby, Phys. Rev

3J

^''M. E. Fine, ASTM Special Technical Publication No. 129,

1952, p. 48 (unpublished).

^^V. H. Robinson, S. H. Carpenter, and J. L. Tallon, J. Appl.

Phys. 45, 1975 (1974).

'''The use in this paper of trade names of specific products

is essential to a proper understanding of the work -presented.

Their use in no way implies any approval, endorsement,

on recommendation bv NBS. ^V.J. Matheson, J. Phys. E 4, 796(1971). '^C.W. Fowlkes, P. E. Angerhofer, P.N. Newton, and

A. F. Clark, National Bureau of Standards Report No. NBSIR 73-349, 1973 (unpublished). "a. P. Gerk, Rev. Sci, Instrum. 43, 1786(1972).

"d. L. Johnson (private communication).

•'^A.F. Clark, V. D. Arp, W. F. Weston, and J. G. Hust, National Bureau of Standards Report No. NBSIR 74-399, 1974 (unpublished).

""^W. N. Robinson and A. Edgar, IEEE Trans. Sonics LTtrason.

21, 98 (1974).

'"W, Voigt, Lehrhych der Kristallphysik (Teubner, Berlin,

_ 1928), pp. 746-7.

'^.P. Varshni, Phys. Rev. B 2, 3952 1 1970). "'E.R. Naimon, W. F. Weston, and H.M. Ledbetter, Cryogenics 14, 721 (1973).

^'L. D. Landau and E.M. Lifshitz, Theory of Elasticity

(Pergamon, London, 1959), p. 17. "r. F.S. Hearmon, Rev. Mod. Phys. 18, 409(1946).

"^D. H. Chung and W. R. Buessem, Anisotropy in SingleCrystal Compounds , edited by F.W. Vahldiek and S. A, Mersol (Plenum, New York, 1968), Vol. 2, p. 217.

4465 J. Appl. Phys., Vol. 46, No. 10, October 1975

85

W.F.Weston 4465

SEMI-ANNUAL REPORT ON MATERIALS RESEARCH

IN SUPPORT OF SUPERCONDUCTING MACHINERY

FRACTURE MECHANICS PARAMETERS OF ENGINEERING

MATERIALS AT CRYOGENIC TEMPERATURES

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

Cryogenics Division

Institute for Basic Standards

National Bureau of Standards

Boulder, Colorado 80302

April 1976
Summary: Fracture Mechanics Parameters of Engineering

Materials at Cryogenic Temperatures

This section of the report contains two manuscripts. Fracture mechanics design data

are presented for a nickel base superalloy, Inconel 718, and an aluminum base alloy,

2219-T6. The plane strain fracture toughness parameters (Kj^ and Jjc), tensile, and fatigue

crack growth rates (da/dN) were measured using state-of-the-art test procedures. The

results are summarized as follows:

(1) The tensile, fatigue, and fracture properties of Inconel 718 alloy are nearly

constant or slightly improved at decreasing temperatures between 295 and 4 K. The fracr

ture toughness (Kjc) at 4 K is 16% higher than the room temperature value of 96.3 MPa-m'/2.

This alloy is useful for cryogenic structural applications requiring an exceptional yield

strength (1.172 GPa at 295 K) and moderate fracture toughness.

(2) J-integral resistance curves for three specimen thicknesses and valid (according

to ASTM Method E 399) Kjc values at 76 K are reported for aluminum alloy 2219. The

J-integral values were independent of thickness at small crack extensions, but at substantial crack extensions the values for the thin specimens were larger than those for the

thick specimens. The measured Jj^, values were less than those calculated from the measured Kt values. The reason for this discrepancy was that crack extension occurred

before tne Kj^, measurement point was reached.

[

I

88
LOW TEMPERATURE EFFECTS ON THE FRACTURE BEHAVIOR OF INCONEL 718

R. L. Tobler

Cryogenics Division

Institute for Basic Standards

Boulder, Colorado 80302

ABSTRACT

The mechanical properties of a solution treated and double aged

Inconel 718 forging were studied to assess its utility at temperatures

in the ambient- to-cryogenic range. Uniaxial tensile property measurements using unnotched specimens at decreasing temperatures between

295 and 4 K show that yield and ultimate strengths increase by 20%

and 29%, respectively, while ductility remains virtually constant.

Fracture mechanics tests using 2.54 cm-thick compact specimens revealed

that the fatigue crack growth resistance of this alloy improves slightly

at extreme cryogenic temperatures, and its plane strain fracture

1/2 1/2 toughness, K^^, increases from 96.3 MPa«m at 295 K to 112.3 MPa«m

at 4 K. These results are compared with similar data for Inconel 750

alloys.

Key words: Fatigue; fracture; low temperature tests; mechanical properties; nickel alloys

superalloys

.

Contribution of NBS, not subject to copyright.

89
INTRODUCTION

A precipitation hardenable nickel base superalloy known as Inconel 718 was first

developed for elevated temperature service. Then, by virtue of its austenitic structure,

high strength, and weldability, the alloy found applications in the ambient-to-cryogenic

temperature region as well. Current applications at temperatures as low as 20 K include

liquid oxygen and liquid hydrogen pressure vessels for spacecraft. Emerging applications

related to superconducting machinery (torque tubes, damper shields) will probably further

extend the utility of Inconel 718 to temperatures approaching absolute zero.

A need for reliable mechanical property data accompanies the expanded use of this alloy.

Conventional tensile properties for the temperature range 295 to 4 K [1-14] are insufficient

to ctiaracterize the load-carrying capability of large structures subject to fatigue. As demonstrated by this study, relatively brittle fractures may occur when fatigue cracks are located in thick sections of Inconel 718, despite the fact that this material exhibits

appreciable ductility in the unnotched condition.

The existing fatigue crack growth and fracture toughness data [12-17] relate predominantly to thin sections and non-standard specimen geometries, or concentrate only on room

temperature properties. The present study describes the mechanical behavior of Inconel 718

at several temperatures including 295, 195, 76, and 4 K. In addition to conventional tensile

property determinations, fatigue crack growth resistance was evaluated, and the ASTM standard

method of test for plane strain fracture toughness of metallic materials (ASTM E-399-74) [18] was applied to generate valid K^.^ data for this alloy.

MATERIAL

The Inconel 718 alloy was purchased in the form of a forged 7.7 cm x 7.7 cm square bar

that had been annealed at 1200 K for Ih. As quoted from the mill sheet, the chemical

composition in weight percent is: 53.4 Ni, 18.5 Fe, 18.1 Cr, 5.24 Nb + Ta, 2.95 Mo, 0.90 Ti,

0.49 Al, 0.17 Co, 0.10 Mn, 0.10 Cu, 0.10 Si, 0.05 C, 0.01 P, 0.003 B, 0.002 S. All test specimens were machined prior to precipitation hardening, which followed the standard heat

treatment

1) Solution treatment (1256 K for 3/4h, air cool);

and 2) Double age (992 K for 8h, furnace cool to 894 K, hold lOh, and air cool).

PROCEDURES

Tensile

Tensile tests were conducted at room, liquid nitrogen, and liquid helium temperatures at

a crosshead velocity of 0.008 mm per second. Reed [19] described the 44.5 kN screw-driven

machine, cryostat, and low temperature techniques used here. The tensile specimens had a

Tradenames are used in this report for clarity and in no way imply endorsement or recommendation by NBS.

90
reduced section 3.8 cm long and 0.51 cm in diameter. The tensile axis was transverse to the

forging axis, such that the fracture plane orientation matched that of the compact specimens

described below. The outputs of a commercial load cell and clip-on strain gage extensometer

were recorded and analyzed as outlined in ASTM Method E8-72 [20]

.

Fracture

The compact specimens used for fatigue and fracture tests had a thickness, B, of

2.54 cm and a width, W, of 5.08 cm. These specimens were machined in the TS orientation [18],

with their loading axis and crack propagation directions transverse to the original forging

axis. Other specimen features were in accordance with the E-399-74 method, except that a modified notch was incorporated. As shown in Figure 1, the modified notch permits clip gage

attachment at knife edges in the loadline.

Room temperature tests were conducted using a 100 kN servo-hydraulic test machine and

cryostat in ambient air. Low temperature environments were achieved by immersing the load

frame, specimen, and clip gage in alcohol and dry ice (195 K) , liquid nitrogen (76 K) or liquid heliiim (4 K) . The clip gage satisfied E--399-74 linearity requirements at each temperature. A detailed description of the low temperature apparatus and techniques was given by

Fowlkes and Tobler [21].

Fracture toughness specimens were precracked at their K test temperatures and loaded

1/2

to fracture at 1 MPa*m per second. Following the E-399-74 method, K^^ was calculated from

the solution for compact specimens:

= Pq b"^ w"^/^ [f(a/W)] (1)

where

f(a/W) = 29.6 (a/W)-'"''^ - 185.5 (a/W)"^^^ + 655.7 (a/W)^''^

(2) - 1017.0 (a/W)''''^ + 638.9 (a/W)^''^

In these equations, is the load defined by the secant offset procedure [18], and a is

the average of three crack length measurements at 25, 50, and 75% of specimen thickness.

Fatigue

Fatigue crack growth rates were measured during the precracking of K^^ specimens where

the maximum stress intensity factor never exceeded 0.6 K^^. Additional specimens were tested

to obtain growth rates at higher stress intensities. Although most specimens were 2.54 cm thick (W/B = 2), four tests of a 0.508 cm thickness (W/B = 10) were also included. The

fatigue operations were conducted using controlled load, a sinusoidal load cycle, and a cycle

frequency of 20 Hz; the ratio, R, of minimum/maximum load was 0.1.

Crack growth was monitored indirectly by compliance measurements, as previously

described [21]. Briefly, a correlation between crack length and specimen deflection per unit

load (6/P) was determined by measurements on fractured specimens. The fatigue tests were

91
interrupted periodically to record 6/P, from which crack lengths could be inferred and plotted

as a function of fatigue cycles, N. A computer program was used to fit the a-versus-N curves with a third order polynomial, to differentiate the curves to obtain da/dN values, and to calculate the associated stress intensity factor ranges, AK, from the peak fatigue loads:

AK = (P -P^,.„) B"^ w"^/2 [f(a/W)] (3) max mm

RESULTS

Tensile

Table 1 lists the yield and ultimate strengths, elongation, and reduction in area measurements for Inconel 718 at 295, 76, and 4 K. These data are plotted in Figures 2-4, along

with data for other Inconel 718 forgings in the 1256 K solution treated and double aged

condition. The temperature dependences of these tensile parameters illustrate what are considered to be classical trends for face centered cubic metals and alloys [22-24]:

ductility remains nearly constant, there is a moderate increase of yield strength, and a larger increase of ultimate strength. The yield strength of the present alloy ranges from

1.172 to 1.408 GPa over the temperature interval investigated, showing good agreement with

other results [9-13],

On the other hand, the heat-to-heat variations of ductility appear to be sizable. With

reference to Figure 4, reduction in area measurements at room temperature range from 14 to 30%, but the values for our alloy are intermediate at 19%.

Handbooks [1-3] contain other low temperature tensile data for thin sheet and other

Inconel 718 stock having thermal or mechanical treatments differing from the 1256 K solutioning and double aging treatment studied here. The comparisons above are sufficient to demonstrate that the present is representative of commercially available forgings.

Fracture

The load-deflection behavior and fracture surface of a compact specimen tested at 4 K

are shown in Figure 5. "Type I" test records [18] with slight nonlinearity prior to maximum

load were observed at each temperature. The fracture surfaces displayed no shear portions

but were flat and granular in appearance, and uninfluenced by test temperature.

The K calculations are itemized in Table 2. The 2.54 cm-thick specimens tested here

/k \2 amply satisfy the E-399-74 criterion for thickness, B >^ 2.5 { Ic j , which requires a minimum

thickness of 1.9 cm for this alloy at room temperature. Other^lnformation pertinent to assessing the validity of measurements are included in Table 2. In three tests, a/W

ic

exceeded the 0.45 to 0.55 range. However, these three deviations are minor, and any effects

on results are considered negligible compared to the degree of scatter among specimens.

The average K value at each temperature is plotted in Figure 6. The data are mildly

1/2 temperature dependent, showing a 16% increase from 96.3 to 112.3 MPa'm as temperature is

lowered from 295 to 4 K. Logsdon, et al. [12] reported a 14% increase for compact specimens

of an Inconel 718 alloy tested at 295 and 4 K. Note that their values are lower (70 to

Ic

92
1/2

80 MPa'm ) compared to the present results; They also reported lower tensile ductility, as shown in Figures 3 and 4. Note also that the temperature dependence of fracture toughness

for Inconel 718 is opposite to that for Inconel 750 aZloys: as temperature is lowered,

Inconel 750 shows moderate decreases in K^^ amounting to 8% between 295 and 76 K [26], and

11% between 295 and 4 K [25].

The literature [13-16] contains other fracture toughness data for Inconel 718 which were

not obtained by the E-399-74 method. Those measurements were derived from tests of centernotched, single-edge-notched, or surface-flawed specimens. To their advantage, surface-flawed

specimens simulate a flaw type that commonly occurs in pressure vessels, but the fracture

toughness results are directly applicable only when the service conditions match the specifics

of specimen flaw shape and thickness. Hall's experimental results indicate that surface

flawed specimens produce high apparent fracture toughness values as compared with the true K^^

values from standard compact or bend specimens [27]. Further difficulty arises from the fact

that the Inconel 718 specimens tested by Witzell [13], Pettit, Feddersen, and Mindin [14] and

Taylor [16], are too thin (1.6 to 3.5 mm) to avoid size effects according to the E-399-74

criterion. For these reasons, such data are of limited value to designers. On the other hand,

the data of Figure 6 are independent of specimen geometry and represent true material

constants which are generally applicable in design.

Fatigue Crack Growth Rates

As demonstrated by Paris and Erdogan [28], fatigue crack growth rates can usually be

described as power-law functions of the stress intensity factor range, AK: i - w

where C and n are material constants that depend on environment and test variables.

Logarithmic plots of da/dN-versus-AK measurements that are in agreement with Eq. (4) reveal

linear trends from which n and C can be determined as the slope and ordinate intercept at AK = 1, respectively.

The fatigue crack growth rates of Inconel 718 at 295, 195, 76, and 4 K are plotted on logarithmic coordinates in Figure 7. The data define a scatterband approximately 7 to 12 1/2 MPa'm wide. At equivalent AK values, the rates at cryogenic temperatures are slightly

lower than at room temperature. However, the improvement in fatigue crack propagation

resistance is modest, and barely distinguishable beyond the scatter of replicate tests unless

the temperature differential exceeds 150 to 200 kelvins.

Linear segments representing two Paris equations were chosen to approximate the upper

(295 K) and lower (4 K) extremes of data in Figure 7. Thus, at 295 K: ^ = 8 X 10"^^ (AK)^ (5)

and at 4 K:

93
da

dN

= 4.8 X 10

,-11 (AK)4 (6)

where AK is in MPa'm'1/2

and da/dN is in mm/cycle. The exponent, n = 4, which some theories

predict [29], provides a good approximation for these experimental results.

Figures 8 and 9 compare the fatigue crack propagation resistance of Inconel 718 with the

available results for other nickel base alloys and stainless steels which could be considered

material competitors for various applications. Note that the room temperature crack growth

rates reported by Shahinian, Watson, and Smith [17] are in close agreement with the present

data for Inconel 718. As shown in Figures 8 and 9, the fatigue crack propagation resistance

of Inconel 718 is intermediate compared with other austenitic alloys.

The low temperature tensile properties of this Inconel 718 alloy were equivalent or superior to its properties at room temperature. Other studies of Inconel 718 alloys have

shown that, as the temperature decreases to 4 K, Young's modulus increases by 6% [31], fatigue

strength at 10 cycles increases by approximately 75% [8] , and Charpy impact values show

little change [11]. Since none of these parameters are deleteriously affected, the reliability

of Inconel 718 structures will not be decreased at extreme cryogenic temperatures.

The attractive feature of a high yield strength which increases at cryogenic temperatures

without sacrifice of tensile ductility justifies Inconel 718 's selection for numerous

applications. Nevertheless, this alloy is susceptible to brittle fracture in thick sections

containing flaws. The fracture toughness at ambient and cryogenic temperatures is useful but

moderate; therefore, accurate fracture mechanics evaluations seem necessary to assure

structural safety and efficient design.

In the fracture mechanics approach, K is proportional to the magnitude of elastic

stresses at the crack tip. Upon reaching a critical value, K^^, the material fractures

catastrophically , without significant plastic deformation. Using handbook [32] equations for

various component geometries of known dimensions and loadings, the maximum flaw size that a structure will tolerate can be calculated based on the material's K^^ value. Fatigue crack

growth rates enable a prediction of the total number of load cycles required to propagate

an initial crack to critical proportions. Thus, the da/dN and K^^ data of this report can

be used to assess the safe operating lifetimes of Inconel 718 structures.

The significant finding that Inconel 718 's da/dN and K^^ parameters are nearly invariant

at temperatures below 295 K should facilitate fracture mechanics analyses. For example, the

performance of a component of this material at cryogenic temperatures could be conservatively

predicted using room temperature data obtained from quality control tests of production

material. Similarly, for convenience and economy, cryogenic structures might be proof-tested

at room temperature.

An important question regarding alloy selection is whether Inconel 718 might supersede

Inconel 750 as a construction material for some applications in superconducting generators.

DISCUSSION

94
Inconel 750 and 718 alloys possess a high resistance to fatigue crack propagation which seems

to be a characteristic of austenitic alloys in general [30]. Figures 8-9 compared the fatigue

crack growth rates of Inconel 750 and 718 with those reported for several austenitic stainless

steels. Dismissing minor differences in the rates for Inconel 750 and 718, attention centers

on yield strength and fracture toughness comparisons. At 4 K, the yield strengths of solution

treated and double aged Inconel 750 alloys range from 0.736 to 1.194 GPa, depending on manufacturing process and heat treatment variations [25] ; K values for the same heats range

1/2 from 102 to 239 MPa*m . Based on the present study, Inconel 718 offers a higher yield

1/2 strength, 1.408 GPa, but its K^^ value of 112 MPa*m is on the low side of the range for

Inconel 750 alloys. This trade-off between yield strength and fracture toughness will

determine alloy selection relative to the specific requirements of each application.

CONCLUSIONS

Uniaxial tensile properties and fracture mechanics data were obtained for a precipitation

hardened Inconel 718 forging at temperatures in the ambient-to-cryogenic range. The temperature dependences of the mechanical properties of this alloy may be considered typical,

qualitatively, of austenitic alloys in general. At 4 K, the yield strength of this Inconel

718 alloy was 20% greater than the room temperature value of 1.172 GPa. The K values

1/2 showed that fracture toughness varies from 96.3 to 112.3 MPa*m between 295 and 4 K, whereas

tensile ductility and fatigue crack growth resistance remain nearly constant. Based on these

data trends, the load-carrying capability of flawed or unflawed Inconel 718 structures is not

expected to decrease at cryogenic temperatures.

ACKNOWLEDGEMENTS

This program was sponsored by the Advanced Research Projects Agency of the U.S.

Department of Defense. The contributions of Dr. R. P. Mikesell and R. L. Durcholz in

conducting the tensile tests are gratefully acknowledged.

95
REFERENCES

1. Handbook on Materials for Superconducting Machinery, MCIC-HB-04 (Battelle Columbus

Laboratories, 1974).

2. Cryogenic Materials Data Handbook, AFML-TDR-64-280 (Martin Marietta Co., 1964).

3. Materials Data Handbook: Inconel Alloy 718, NASA-CR-123774 (Western Res. and Dev. Inc.,

1972)

4. Inconel Alloy 718, Bulletin T-39 (Huntington Alloy Products Division, International

Nickel Co. , 1973)

5. Wagner, H. J., Hall, A. M. , Physical Metallurgy of Alloy 718, Defense Metals Information

Center Report 217 (Battelle Memorial Institute, 1968).

6. Christian, J. L., in Advances in Cryogenic Engineering 12^ (1962) 520.

7. Nachtigall, A. J., in Properties of Materials for Liquefied Natural Gas Tankage, ASTM

STP 579 (Amer. Soc. Test. Mater., Philadelphia, 1975) 378.

8. Nachtigall, A. J., Klima, S. J., Freche, J. C, J. Materials, JMLSA _3 (1968) 425.

9. Inouye, F. T., Hunt, V., Janser, G. R. , Frick, V., NASA Tech. Rept. CR-788 (Aerojet

General Corp. , 1967)

10. Morgan, W. R. , NASA Internal Note IN-P and VE-M-66-2 (George C. Marshall Space Flight

Center, 1966).

11. Malin, C. 0., NASA Tech. Brief MFS-18244 (George C. Marshall Space Flight Center, 1965).

12. Logsdon, W. A., Kossowsky, R. , Daniel, M. R. , Wells, J. M. , in Materials Research for

Superconducting Machinery IV, Reed, R. P., Clark, A. F., and van Reuth, E. C, Eds.,

Semi-Annual Tech. Rept., October 1975 (Advanced Research Projects Agency, 1975).

13. Witzell, W. E., Tech. Rept. AFML-TR-67-257 (General Dynamic s-Convair, 1967).

14. Pettit, D. E., Feddersen, C. E. , Mindin, H. , NASA CR 101942 (Battelle Memorial Institute,

1969).

15. Forman, R. G. , NASA Tech. Note TN D-7665 (National Aeronautics and Space Administration,

1974).

16. Taylor, J. L., in Proceedings of the Seventh Symposium on Nondestructive Evaluation of

Components and Materials in Aerospace (Amer. Soc. Nondestructive Testing, San Antonio,

1969) 420.

17. Shahinian, P., Watson, H. E. , Smith, H. H. , J. Materials, JMLSA 7^ (1972) 527.

18. Annual Book of ASTM Standards, Part 10 (Amer. Soc. Test. Mater. Philadelphia, 1974) 432.

19. Reed, R. P., in Advances in Cryogenic Engineering 2_ (1962) 448.

20. Annual Book of ASTM Standards, Part 31 (Amer. Soc. Test. Mater., Philadelphia, 1973) 220.

21. Fowlkes, C. W. , Tobler, R. L. , Eng. Fract. Mechs., to be published.

22. McLean, D., Mechanical Properties of Metals (Wiley, New York, 1962).

23. Conrad, H. , in High Strength Materials (Wiley, New York, 1964).

24= Wigley, D. A., Mechanical Properties of Materials at Low Temperatures (Plenum, New York,

1971).

25. Logsdon, W. A., in International Cryogenic Materials Conference Proceedings (Plenum,

New York, 1976) to be published.

96
26. Tobler, R. L. , Mikesell, R. P., Durcholz, R. L. , Reed, R. P., in Materials Research for

Superconducting Machinery II, Clark, A. F., Reed, R. P., and van Reuth, E. C, Eds.,

Semi-Annual Tech. Rept., October 1974 (Advanced Research Projects Agency, 1974).

27. Hall, L. R., in Fracture Toughness Testing at Cryogenic Temperatures, ASTM STP 496

(Amer. Soc. lest. >Iater., Philadelphia, 1970) 40.

28. Paris, P. C, Erdogan, F., J. Basic Engrg., Trans. ASME Series D, 85 (1963) 528.

29. Johnson, H. H. , Paris, P. C, Eng. Fract. Mech. 1_ (1968) 3.

30. Tobler, R. L. , Reed, R. P., in International Cryogenic Materials Conference Proceedings

(Plenum, New York, 1976) to be published.

31. Weston, W. F., Ledbetter, H. M. , Mater. Sci. Eng. 20_ (1975) 287.

32. Tada, H. , Paris, P. C, Irwin, G. R. , The Stress Analysis of Cracks Handbook (Del

Research Corp., Hellertown, PA, 1973).

97
Table 1. Tensile property results for Inconel 718.

Temperature 0. 2% Yield Strength Ultimate Strength Elongation Reduction in Area

K GPa (ksi) GPa (ksi) % . %

295 1.165 1.391 16.1 19.5

1.396 14.3 18.4

1.178 1.425 15.8 16.8

Avg 1.172 (170.0) 1.404 (203.6) 15.4 18.2

76 1.326 1.649 20.2 20.6

1.359 1.649 20.9 18.9

Avg 1.342 (197.4) 1.649 (239.2) 20.6 19.8

4 1.371 1.802 21.7 24.0

1.445 1.830 19.5 16.3

1.408 (204.2) 1.816 (263.4) 20.6 20.2

98
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99
LIST OF FIGURES

Figure 1. Compact specimen used for fatigue and fracture tests.

Figure 2. Yield and ultimate tensile strengths of Inconel 718 forgings in the solution

treated (1225 to 1256 K) and double aged condition.

Figure 3. Tensile elongation of Inconel 718 forgings in the solution treated and

double aged condition.

Figure 4. Reduction in area of Inconel 718 forgings in the solution treated and double

aged condition.

Figure 5. Fracture toughness test record for Inconel 718, showing fracture surface

appearance.

Figure 6. Temperature dependence of fracture toughness for two Inconel 718 alloys.

Figure 7. Fatigue crack propagation rates for Inconel 718.

Figure 8. Comparison of fatigue crack propagation rates of Inconel alloys 718 and 750.

Figure 9. Comparison of fatigue crack propagation rates of Inconel 718, Inconel 750, and

austenitic stainless steels at 4 K.

100
3.1 cm

6.1 cm 6a

y^' 1.27cm dia.

1.52 cm

">H^ I 3.05cm

.79 cm dia.'^^

I j

W=5.08cm^

6.35cm ^

THICKNESS, B = 2.54 cm; W/B = 2

Figure 1. Compact specimen used for fatigue and fracture tests.

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105
106
5

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STRESS INTENSITY RANGE, K, MPa-m'^^

Figure 7. Fatigue crack propagation rates for Inconel 718.

107
5

E 10

E

Z

D

-3

5

o

<

<

10

-5

Inconel 718, This Study

4<T <295K

Inconel 718, Shahinian (17)

•—• T = 295K

Inconei 750 alloys

Logsdon (25)

4<T<300K

10 100

STRESS INTENSITY RANGE, aK, MPa m

Figure 8. Comparison of fatigue crack propagation rates

of Inconel alloys 718 and 750.

108
5

T = 4K

1,0-'

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0 AISI 304

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® Inconel 750

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Inconel 718

(This Study)

J I I I I I I i

5 10 50 100

STRESS INTENSITY RANGE, aK, MPa m^/^

Figure 9. Comparison of fatigue crack propagation rates of Inconel

718, Inconel 750, and austenitic stainless steels at 4 K.
EFFECTS OF SPECIMEN THICKNESS ON FRACTURE

TOUGHNESS OF SOME ALUMINUM ALLOYS*

D. T. Read"^ and R. P. Reed

Cryogenics Division

Institute for Basic Standards

National Bureau of Standards

Boulder, CO 80302

This research was partically supported by the Advanced Research Projects

Agency of the U. S. Department of Defense; it is a contribution of NBS,

not subject to copyright.

NRC-NBS Postdoctoral Research Associate, 1974-1975.

110
EFFECTS OF SPECIMEN THICKNESS ON FRACTURE

TOUGHNESS OF SOME ALUMINUM ALLOYS

J-integral resistance curves for three specimen thicknesses and valid

(according to ASTM Method E 399) K^^ values at 76 K are reported for aluminum

alloy 2219. The J-integral values were independent of thickness at small

crack extensions, but at substantial crack extensions the values for the

thin specimens were larger than those for the thick specimens. The measured

J, values were less than those calculated from the measured values. The

Ic Ic reason for this discrepancy was that crack extension occurred before the Kj^

measurement point was reached.

Key words: Aluminum alloys; critical stress intensity, cryogenic

temperature; fracture toughness; J-integral; sub-critical crack

extension.

111
1. Introduction

The generally accepted measure of the fracture toughness of a metal

is the plane strain fracture toughness, Kj^, as measured by ASTM Method

E 399-74 (Ref. 1). However, the specimen size required for a valid

measurement of Kj^ may be so great that tests to obtain Kj^ may be too

difficult or too expensive to perform (Ref. 2). An alternative method

of measuring plane strain fracture toughness known as the J-integral

method (Refs. 3, 4) is being widely considered because of its less stringent requirements on specimen size. The results described in this

report are of use in evaluating the J-integral method of fracture toughness testing.

The parameter characteristic of fracture toughness which is measured in a J-integral test is Jj^, the value of the J-integral at the

onset of crack extension. In the linear elastic case, Kj^ and J^^ are

related as (Ref. 3)

J, = iic (1 - v^) (1)

E

where E is Young's modulus and v is Poisscn's ratio of the material.

The J-integral test can be traced back to the Griffith criterion

(Refs. 5, 6) for crack extension and to the analytical work of Eshelby

(Ref. 7). However, Rice (Ref. 8) and Begley and Landes (Refs. 3, 4)

have been mainly responsible for the formulation of the current concept of the J-integral and its uses in fracture toughness testing.

Landes and Begley (Ref. 4) have recently proposed a standard test

method for determining J^ .

112
There are two major obstacles to the successful implementation of

the J-integral test method. The proposed test method is an attempt to

resolve the first of these, namely, the choice of a measurement point.

The basic procedure in a J-integral test is to determine experimentally

a resistance curve consisting of a set of measured J-integral versus

crack extension (J versus Aa) data points, and then to determine Jj^ from

this curve according to a rule which specifies the measurement point. In

the proposed J-integral test method, the measurement point is chosen at a

knee in the J-Aa curve, and thus the shape of this curve is critical to

the measurement. It is asserted that if Jj^ is chosen properly, Eq.(l)

holds, and a measurement of Jj^ is equivalent to a measurement of Kj^.

The second major problem with the J-integral method is that the validity

of Eq.(l) for useful structural materials has not been demonstrated theoretically (Ref. 9). The attempted extension of the equivalence of Jj^ and

Kj^ from the linear elastic to the elastic plastic case has been supported

by appeal to experimental results. There is evidence that for some materials Jj^ and are equivalent (Ref. 4). However, J^^ and Kj^ have

been found not to be related by Eq. (1) in some other materials (Ref. 10).

The objectives of the present study were to determine experimentally:

(1) Jj^ as a function of specimen thickness for an aluminum alloy;

(2) the relationship between Jj^ and Kj^, when K^^ is determined

according to ASTM Method E 399;

(3) the effect of specimen thickness on the shape of the J-Aa curve.

Information of this kind is quite scarce, and is necessary before the

J-integral method can be standardized as a reliable test method.

113
2. Materials

Two aluminum alloys, 2014-T652 and 2219-T87, were used as specimen materials. The chemical compositions of these alloys are given in

Table 1. The 2014-T652 designation indicates that this material has been

heat-treated, artificially aged, and stress relieved by compressing. The

2219-T87 designation indicates that this material has been solution heattreated, cold-worked, and then artificially aged. After polishing and

etching, the grains in both alloys were seen to be severely elongated.

A section cut perpendicular to the rolling direction shows some grains about

.02 cm (.008 inch) in diameter and some larger grains of various sizes

in the 2014 alloy, while the same treatment reveals grains about .03 cm

(.012 inch) in diameter in the 2219 alloy. On sections cut parallel to

the rolling direction, only grain boundaries approximately parallel to

the rolling direction, spaced .02 cm were seen in the 2014 alloy. The

2014-T652 and 2219-T87 alloys had Rockwell hardness of B 80.6 and B 81.1,

respectively.

The yield strength a^, ultimate tensile strength cr^^^, Young's

modulus E, Poisson's ratio v, and plane strain fracture toughness Kj^ values

obtained for the two alloys tested are listed in Tables 2 and 3 (Ref. 11).

For measurements of Kj^ and J, the 2014 alloy specimens were machined in

the L-T orientation, and the 2219 specimens were machined in the T-L orientation. For comparison, handbook (Ref. 12) values of Kj^ are also shown.

The unusually high Kj^ of the 2014 alloy specimens is probably due to

nonuniformi ties in the specimen material produced during forging. Metallographic evidence for such nonuniformi ties is discussed below. Originally,

the 2014 alloy was chosen for its availability and the convenience of

machining and testing the specimens. The 2014 alloy specimens were cut

114
from a 7.6 cm (3 inch) thick forged plate.

The J-_a data for the 2014 alloy summarized in Table 4 exhibited an

uneven dependence of J upon la. This scatter v;as attributed to variation of properties among the specimens. The specimens used in the tests

were examined metallographical ly. On sections cut perpendicular to the

rolling direction, some relatively large grains several millimeters in

diameter were observed along with the small grains .02 cm in diameter.

The fraction of the surface which was made up of small grains was recorded

in Table 4 for the 2014 alloy specimens. Both the raw J-la data and the

data classified by grain size for the 2014 alloy exhibited so much scatter that they were useless for the purposes of the present study. For

this reason, a plate 3.8 cm (1.5 inch) thick of aluminum alloy 2219-T87

was obtained and tested. The remainder of this report is concerned with

the results of measurements made on the 2219 alloy.

3. Procedures

The specimen geometry used was a modified compact specimen, as

shown in Fig. 1. The normal dimensions of all the specimens tested are

listed in Table 5. Tests v/ere conducted in an electrohydraul ic testing

apparatus equipped with a stand-off and dewar for cryogenic service which

has been described previously (Ref. 13), The specimens were mounted in

the loading fixtures and cooled to 76 K by immersion in liquid nitrogen.

Each specimen was precracked by loading in tension-tension fatigue at

20 Hz at a load which in most cases was 60': or less of the quasi static

load used in the J test. The precracking was continued to a nominal crack

length of a/W = .60. Typically the precrack reached the desired length

of about .5 cm (.2 inch) after approximately 60,000 fatigue cycles.

After precracking, each specimen was loaded quasi statical ly to a stress

115
equal to or less than the maximum load used in the precracking procedure,

and the load-displacement line was recorded. This line was extended to

high loads, and a secant line deviating from this anticipated load line

by a preselected amount was also drawn. After these preliminary steps,

the quasistatic crack extension was accomplished by loading the specimen

until the actual load-displacement curve intersected the pre-drawn

secant line, and immediately unloaded. Then the specimen was refatigued,

beginning with a load of about half the maximum load reached during the

measurement of the load-displacement curve. After the refatigue crack

had grown to several millimeters in length, the specimen was pulled apart

so that the fracture surfaces were accessible for examination. The crack

growth produced during the static load cycle was visible as a dark, rough

band between the two bright, relatively smooth, fatigue surfaces.

Each value of the J-integral was computed from the load-displacement

record obtained during the static load cycle, using the procedure recommended by Merkle and Corten (Ref 14). This procedure results in a value

of the J-integral of

J

= AA, (2)

Bb

where A is the area under the load displacement curve, B is the specimen

thickness, b is the specimen ligament remaining after precracking, and

X is a function of the crack length and the nonlinearity of the loaddisplacement record. The value of A was approximately 2.4 in all the tests

performed in this study.

In the proposed standard method for determining J, (Ref. 4) the

Ic

crack extension Aa is defined to include both the stretch zone and the

zone of material separation, if any. If the crack extension consists

116
only of a stretch zone, the J-Z.a graph is expected to fall along the

line given by

J = 2 a.f 1 ow

Aa (3)

where aflow

is the arithmetic mean of the yield strength and the ultimate tensile strength. J^^ is defined as the value of J at the intersection of the line given by Eq.(3) with the back-extrapolation of the

J-Aa curve from the region where material separation does take place.

According to the proposed method, the back-extrapolated J-A curve is

expected to have a slope different from that of the line given by Eq.(3),

so that the measured J-Aa curve ought to have a relatively sharp knee

in it.

Because of the relatively high value of the yield and ultimate

tensile stresses of the 2219-T87 alloy, J^^ as defined by this proposed

standard occurs at a crack extension of less than ,002 cm. Due to the

difficulty of measuring extremely small values of crack extensions, and

the apparent unreliability of extrapolating to zero crack extension, the

initiation of crack extension was taken as the point of .005 cm crack

extension for purposes of the present study. This point was determined

by interpolation or extrapolation on graphs of J-Aa. The J^^ values

obtained in this way were higher than those which could have been obtained

by extrapolation according to the proposed standard, but they were

less uncertain.

The J-Aa data for aluminum alloy 2219 are listed in Table 6 and

plotted in Fig. 2. The J^^ values obtained from the J-Aa curves are

listed in Table 7. The values of J^ range from 9.5 to 10.3 kJ • m~

4. Results and Discussion

117
over a tenfold variation in specimen thickness. Thus, J^^ is indeed a

thickness-independent measure of the fracture resistance of the material.

Using the measured value of K^^ and Eq. (1), one would expect a Jj^

-2

of 13.6 kJ • m , 38% larger than the values actually found. Thus, for

aluminum alloy 2219, Kj^ and Jj^ are not related by Eq. (1). This finding

was attributed to the difference in the measurement point for the two test

methods; subsequent experimental measurements showed this to be the correct

explanation for the observed discrepancy. The measurement point for determining Jj^ was chosen as the initiation of crack extension, whereas the

measurement point for determining Kj^ was chosen as set forth in ASTM

Method E 399, that is, at 2% apparent crack extension. The apparent crack

extension includes the plastic zone, but the possibility exists that

significant actual crack extension may occur before the Ki^ measurement

point is reached. If this is the case, the values of K^^ obtained may be

characteristic of a somewhat extended crack, and not of crack initiation.

This possibility is widely suspected (Ref. 15) as a cause of discrepancy

between observed values of Kj^ and J^^ as related by Eq. (1), and was used

by Logsdon (Ref. 10) to explain some of his data. But it appears that

this explanation has never been tested experimentally.

To determine the extent of crack extension in a valid K^^ test,

three 3.8 cm (1.5 inch) thick compact tensile specimens, more than twice

as thick as necessary to obtain valid Kj^ data, were tested according to

the J-procedure described in Section 3. In two tests, the specimen was

unloaded prior to the intersection of the load line with the 5% offset

secant, and in the other test the specimen was unloaded just after this

intersection. The records of these tests were analyzed as follows. A

118
line was drawn through the origin of the test record and the point on the

record at which unloading was begun. The percent difference between the

slope of this secant line and the tangent to the initial linear portion

of the load record, referred to as the percent secant deviation, was

recorded. This procedure allowed a correlation of secant deviation with

crack extension to be made. The data are listed in Table 8 and plotted

in Fig. 3. These data clearly show that the critical point in the K^^

test of 2219 aluminum alloy, that is, the point of 5% secant deviation,

occurs at a crack extension of about .064 cm (.025 inch) in specimens

which satisfy the thickness requirements by a wide margin. As a double

check, a similar plot for the 1.42 cm (.56 inch) thick specimens, which

are almost exactly in the minimum thickness required in a Kj^ test, shows

the same characteristic; i.e., crack extension of about .064 cm (.025

inch) before the 5% secant offset is reached. If the Jj^ evaluation

point were chosen not at crack initiation but at .064 cm (.025 inch) crack

extension in these specimens, good agreement could be found between

Kj^ and Jj^. Thus, the discrepancy between the measured values of K^^

and Jj^ related by Eq. (1) is explained. However, because of the observed

dependence of the J integral value at significant crack extensions on

specimen thickness, the critical state for J^^ measurement cannot be chosen

as a point of significant crack extension uch as .064 cm (.025 inch)

because this would defeat the purpose of J-integral testing, which is to

obtain a parameter characteristic of the material tested independent of

specimen thickness. Thus, in this alloy it is impossible to obtain agreement between Kj^ and Jj^ in Eq.(l) that is independent of thickness.

An obvious feature of the J-resistance curves of Fiq. 2 is that the

curves for the three different material thicknesses separate for

119
the larger values of crack extension but converge near zero crack extension. It can be seen for this material that the effect of increasing

specimen thickness is to decrease the J values at sizable values of Aa

while leaving them unchanged, within experimental uncertainties, near

zero crack extension. A similar divergence in thickness was observed

by Griffis and Yoder (Ref. 16) for 2024-T351 aluminum tested at room temperature using a three-point-bend specimen and thus may be a normal

feature of J-integral behavior of ma>ny aluminum alloys. Thickness

effects are discussed further in the Appendix below.

The following conclusions have been drawn from the present study:

(1) In specimens cut from a plate of 2219 aluminum alloy, crack

extension was initiated at a constant value of the J-integral over a ten-fold range of specimen thickness, extending from well above to well

below the minimum specimen thickness given in ASTM Method E 399 for valid

Kj^ measurement. This result supports the hypothesis that Jj^, is a useable measure of material toughness.

(2) and K^^ (K^^ determined according to ASTM Method E 399) are

not equivalent measurements of fracture toughness for 2219 aluminum because

of subcritical crack growth in Kj^ tests; the Jj^ value anticipated from

measurements of K^^ was 38% larger than the measured J^^ value.

(3) As crack extension increased, the J-integral value was significantly larger for thin specimens than for thick ones.

Appendix. Observations on Plane Stress Toughening

The specimen thickness B may be expressed in dimensionl ess form:

5. Concl usions

B Y (Al) =

2

120
The parameter y can be used as an index of the operative fracture mode.

The possible modes are plane strain, mixed, and plane stress. According

to Eq.(Al) and the measured values of the parameters Kj^ and a^, the samples used in the present study fall at y = -72, 1.34, 2.68, and 7.19.

Values of y equal to or greater than 2.5 correspond to the criteria of

ASTM Method E-399 for plane strain testing. Thus, for values of y below

2.5, we may expect plane stress toughening effects, whereas for y above

2.5 plane strain should dominate. A strong increase in apparent sample

toughness is to be expected as the thickness is decreased (Ref. 17).

An experimental search for such a dependence of apparent toughness

on specimen thickness may be made by observing the maximum stress intensity sustained by the samples as a function of thickness. Such stress

intensities may be calculated using the maximum load sustained by the

sample and the crack length present at the initiation of loading. Values

of the maximum stress intensity measured in this manner are denoted K|^.

may be corrected for the plastic zone size as follows. First, the

plastic zone radius is calculated from the uncorrected stress intensity

according to (Ref. IB):

Then the effective crack length a = a + r is used to find the first

eff y

corrected stress intensity factor according to:

r y

(A2)

P f(aeff/W ) iV2

K Ic

(A3)

B W

121
where f( —jj^) is given by Newman (Ref. 19). This corrected stress intensity is used to calculate a new value of r^, which is then used to calculate a new stress intensity, and the iteration continues in this manner.

If this process converges, the final value of r^ is not too large and the

plastic zone correction has been made successfully. An asterisk indicates

a stress intensity value which includes the plastic zone correction. For

instance, K„. is the stress intensity at maximum load corrected for the

plastic zone.

Measured values of K^. and K... are listed in Table A-1 and plotted

versus thickness in Fig. A-1. The classic decrease of this measure of

material toughness with increasing specimen thickness is emphasized by

the plastic zone correction. The curve shown in this figure is given by

where Y was defined in Eq.(Al). This equation was used by Irwin et al

(Ref. 19) to curve-fit some fracture toughness data on aluminum. The

qualitative agreement between this curve and the measured values of K^^^

indicates that the anticipated strong increase in apparent sample toughness with decreasing thickness was indeed found in the present study.

The increase of Kj^^ indicates that plane stress toughening occurs as the

specimen thickness decreases. From these data, we may conclude that the

border between plane strain and mixed modes occurs at about y = 2.5; thus,

the requirements of ASTM Method E-399 are adequate to insure plane strain

conditions, but are not excessive- Due to limitations on the

specimen size, the pure plane stress region was not attained in these

tests. Nevertheless, the data are sufficient to show tendencies of the

difference between plane strain and plane stress modes of fracture.

(A4)

122
References

1. Standard Method of Test for Plane Strain Fracture Toughness of Metallic Materials (Designation E 399-74), 1974 Annual Book of ASTM Standards, Part 10, 1974, pp. 432-451.

2. Kaufman, J. 6. and Kelsey, R. A., in: Properties of Materials for

Liquefied Natural Gas Tankage , ASTM STP 579, American Society for

Testing and Materials, Philadelphia, PA (1975), pp. 138-158.

3. Begley, J. A. and Landes, J. D. , in: Fracture Toughness , Proceedings

of the 1971 National Symposium on Fracture Mechanics, Part II, ASTM

STP 514, American Society for Testing and Materials, Philadelphia,

PA (1972), pp. 1-20.

4. Landes, J. D. and Begley, J. A., in: Fracture Analysis , ASTM STP

560, American Society for Testing and Materials, Philadelphia, PA

(1974), pp. 170-186.

5. Griffith, A. A., Phil. Trans. Royal Soc. 221A , 163 (1920-21);

Reprinted with corrections in Trans. Quart. ASM 6J_, 861 (1968).

6. Griffith, A. A., Proc. First International Congress on Applied Mechanics, Delft (1924) , p. 55.

7. Eshelby, J. D. , "The Force on an Elastic Singularity," Trans. Roy.

Soc. Lond. A244, 87 (1951).

8. Rice, J. R. , in: Fracture , Vol . II , H. Liebowitz, ed., Academic

Press, New York (1968), p. 191.

9. Eftis, J. and Liebowitz, H., Engr. Fracture Mechanics 7_, 101 (1975).

10. Logsdon, A. W. , "Elastic Plastic (Jj^) Fracture Toughness Values:

Their Experimental Determination and Comparison with Conventional

Linear Elastic (Kj^) Fracture Toughness Values for Five Materials,"

Westinghouse Research Laboratories, Pittsburgh, PA, Scientific Paper

74-1E7-FMPWR-P1 (1974).

123
11. The yield strengths, ultimate tensile strengths, and fracture

toughnesses were measured using standard techniques; Young's

modulus and Poisson's ratio for the specimen materials were

measured by an ultrasonic technique as described by D. T. Read

and H. M. Ledbetter, to be published.

12. Matthews, W. T., "Plain Strain Fracture Toughness (Kj^,) Data Handbook for Metals," Army Materials and Mechanics Research Center, AMMRC

MS 73-6, Watertown, MA (1973).

13. Fowlkes, C. W. and Tobler, R. L. , "Fracture Testing and Results

for a Ti-6A1-4V Alloy at Liquid Helium Temperature," to be published in Engineering Fracture Mechanics.

14. Merkle, J. G. and Corten, H. T. , "A J-Integral Analysis for the Compact Specimen, Considering Axial Force as Well as Bending Effects,"

presented at the Pressure Vessels and Piping Conference with Nuclear

Engineering and Mechanical Divisions, American Society Mechanical

Engineers, Paper No. 74-PVP-33 (1974).

15. Tobler, R. , Sha, G. and Begley, J. , personal communication.

16. Griffis, C. A. and Yoder, G. R. , "Application of the J-Integral to

Crack Initiation in a 2024-T351 Aluminum Alloy," Naval Research

Report 7676, (1974), p. 21.

17. Srawley, J. E. and Brown, W. F. , in: Fracture Toughness Testing and

Its Applications , ASTM STP 381, American Society for Testing and

Materials, Philadelphia, PA (1965), pp. 133-196.

18. Weiss, V. and Yukawa, S., in: Fracture Toughness Testing and Its

Appl i cations , ASTM STP 381 , American Society for Testing and Materials, Philadelphia, PA (1965), pp. 1-22.

124
19. Newman, J. C, Jr., in: Fracture Analysis , ASTM STP 560, American

Society for Testing and Materials, Philadelphia, PA (1974), pp. 105-121.

20. Irwin, G. R. , Krafft, J. M. , Paris, P. C. , Wells, A. A., "Basic

Aspects of Crack Growth and Fracture," Naval Research Laboratory

Report No. 5598, 1967, p. 50.

125
LIST OF TABLES

Table 1. Chemical compositions of the alloys, wt. %.

Table 2. Tensile and fracture data for aluminum alloy 2014-T652 in

the L-T orientation.

Table 3. Tensile and fracture data for aluminum alloy 2219-T87' in

the T-L orientation.

Table 4. J-integral data for aluminum alloy 2014-T652 in the T-L

orientation at 76 K.

Table 5. Dimensions of specimens used in the present study.

Table 6. J-integral data for aluminum alloy 2219-T87 in the L-T

orientation at 76 K.

Table 7. Measured and predicted (from Kj^) values of Jj^ for four

specimen thicknesses of aluminum alloy 2219-Ta7 at 76 K.

Table 8. Measured values of crack extension and secant deviation

for two specimen thicknesses of aluminum alloy 2219-T87 at

76 K.

Table Al . Maximum stress intensity and maximum stress intensity cor- rected for the plastic zone for four specimen thicknesses

of aluminum alloy 2219-T87 at 76K.

Table 1. Chemical compositions of the alloys, wt

Cu Fe Mg Mn Si Ti V Zn Zr

2014 4.4 0.5 0.8 0.8

2219 6.4 0.20 <.01 0.26 0.15 0.16 0.12 0.09 0.16

For 2014, nominal composition. For 2219, plasma arc analysis.

126
T(K)

in the L-T orientation.

(MPa)

(ksi)

o ^ (ksi) uts

E(GPa)

E(Ksi)

V

(MPa

Ic

1/2, m )

K^^ (ksi . in^/2)

1/2

(MPa ' m ' ) ic

(handbook value)

1/2 (ksi

ic

Avg

300

405

402

= 403.5

58.8

58.4

Avg = 58.6

476

475

Avg = 475.5

69.1

68.9

Avg = 69.0

75.6

11.0

.341

53.4

46.7

Avg = 50.1

48. 6

42.5

Avg = 45.6

26.4

24

76

442

436

Avg = 439.

64.1

63.3

Avg = 63.7

524

536

Avg = 530.

76.0

77.7

Avg =

Avg =

Avg

76.8

83.3

12.1

.331

49.9

53.4

51.3

51.5

45.4

48.6

46.7

46.9

33.0

30

561

534

Avg = 547.

81.3

77.4

Avg = 79.4

700

663

Avg

Avg

681.5

101.5

96.2

98.8

83.8

12.1

.330

53.4

48.6

(handbook value)

127
dilU 1, L ctu l_ u.i. c U.CI i_ca i'-'i- cixu.

T(K)

(MPa)

(ksl)

a (MPa)

uts

^ ^ (ksl) uts

E(GPa)

E(Msl)

V

(MPa

Ic

1/2- m )

(ksi • m ) ic

1/2

(MPa • m ) ic

(handbook value)

K^^ (ksi • in"^^^)

in the T-L orientation.

300

411

412

Avg = 411.5

59.6

59.7

Avg = 59.65

509

503

Avg = 506.

73.8

73.0

Avg = 73.4

77.4

11.2

.330

31.0

30.6

Avg = 30.8

28.2

27.9

Avg = 28.1

33.0

30.0

14

76 4

490 547

496 514

Avg = 493.

71.1

72.0

Avg = 71.6

630

630

Avg = 630.

91.4

91.3

Avg = 91.35

85.2

12.4

.319

37.2

33.5

35.9

33.9

30.5

Avg = 32.7

Avg

Avg =531.

• 79.4

74.6

Avg = 77.0

734

706

Avg = 720.

106.5

102.4

Avg = 104.4

85.7

12.4

.318

(handbook value)

14

128
Table 4. J-integral data for aluminum alloy 2014-T651 in the T-L orientation at 76 K.

Sample ? BCcm) B(iji) b(cm) b(in) Aa(cm) ia(10 ^ in) J(^) J(.^^y^ % small m in grains

A-6 .508 0.2 3.71 1.46 .oil 4.3 33.6 191.9 10%

A-15 .508 0.2 3.56 1.40 .027 10.5 54.0 308.2 30%

A-1 .508 0.2 3.58 1.41 .036 14.2 29.6 169.2 100%

A-2 .508 0.2 3.62 1.42 .071 27.8 56.0 320.0 50%

A-3 .508 0.2 3.66 1.44 .116 45.2 100.5 573.7 50%

B-8 1.27 0.5 3.61 1.42 .009 3.4 13.4 76.8 100%

B-3 1.27 0.5 3.35 1.32 .009 3.5 26.7 152.3 20%

B-7 1.27 Q.5 3.56 1.40 .010 4.0 17.1 97.8 60%

B-2 1.27 0.5 3.43 1.35 .033 13.0 27.3 156.0 100%

3-6 1.27 0.5 3.48 1.37 .045 17.9 24.9 142.0 100%

B-1 1.27 0.5 3.43 1.35 .054 21.2 34.7 197.9 40%

B-5 1.27 0.5 3.45 1.36 .057 22.5 45.2 264.1 0%

B-4 1.27 0.5 3.53 1.39 .141 55.5 37.5 214.2 - 50%

C-6 2.54 1.0 3.63 1.43 .008 3.1 13.2 75.1 50%

C-1 2.54 1.0 3.45 1.36 .017 6.6 21.1 120.5 100%

C-5 2.54 1.0 3.51 1.38 .018 7.0 30.2 172.7 20%

C-8 2.54 1.0 3.63 1.43 .036 14.1 28.6 162.9 60%

C-3 2.54 1.0 3.68 1.45 .056 22.0 33.0 188.7 25%

C-7 - 2.54 1.0 3.51 1.38 .067 . 26.5 40.3 - 230.3 50%

C-4 2.54 1.0 3.38 1.33 .129 50.9 37.6 215.0 50%

C-2 2.54 1.0 3.30 1.30 .218 86.0 59.2 338.3 100%

- OO
Table 5a. Dimensions of aluminum 2014 specimens used in the present

study. B and W have the meanings indicated in Fig. 1.

B

0.51 cm (0.2 in)

1.27 cm (0.5 in)

2.54 cm (1.0 in)

3.81 cm (1.5 in)

W

7.62 cm (3 in)

7.62 cm (3 in)

7.62 cm (3 in)

7.62 cm (3 in)

4.57 cm (1.8 in)

4.57 cm (1.8 in)

4.57 cm (1.8 in)

4.01 cm (1.58 in)

Number of

specimens

tested

Table 5b. Dimensions of aluminum 2219 specimens used in the present

study. B and W have the meanings indicated in Fig. 1.

B

0 .38 cm ( .15 in)

0.72 cm ( .28 in)

1.44 cm ( .56 in)

3.81 cm (1.5 in)

W

5.08 cm (2 in)

5.08 cm (2 in)

5.08 cm (2 in)

7.62 cm (3 in)

3.05 cm (1.2 in)

3.05 cm (1.2 in)

3.05 cm (1.2 in)

4.19 cm (1.65 in)

Number of

specimens

tested

130
Table 6. J-integral data for aluminum allo^' 2219-T87 in the L-T orientation at 76 K.

Sample # B(cm) B(in)

A-4 .38 .15

A-5 .38 .15

A-3 .38 .15

A-1 .38 .15

A-2 .38 .15

B-4 .71 .28

B-5 • .71 .28

B-3 .71 .28

B-1 .71 .28

B-2 .71 .28

C-3 1.42 .56

C-5 1.42 .56

C-1 1.42 .56

C-2 1.42 .56

C-7 1.42 .56

D-7 3.81 1.5

D-6 3.81 1.5

D-5 3.81 1.5

b(cni) b(in) Aa(cm)

2.03 .80 0

2.04 .80 .013

2.03 .80 .033

2.04 .80 .074

2.02 .80 .130

2.03 .80 .001

2.09 .82 .008

2.06 .81 .024

2.11 .83 .097

2.06 .81 .183

2.08 .82 .002

2.02 .80 .013

2.09 .82 .020

1.99 .78 .076

2.01 .79 .157

3.65 1.44 .009

3.47 1.37 .067

-.51 1.38 .072

-a (10 xn) J(^) J( m in

0 10.0 57.2

5.1 12.6 72.1

12.8 16.7 95.4

29.0 21.8 124.6

51.0 33.7 192.3

0.4 9.6 54.7

3.4 10.9 62.2

9.3 13.1 74.8

38.1 12.5 105.5

71.9 27.3 155.8

0.7 9.0 51.5

5.1 11.7 66.8

8.0 13.0 74.1

30.0 16.6 94.6

62.0 20.2 115.3

3.8 10.0 56.9

26.2 13.4 76.4

28.5 14.8 84.8
Table 7. Measured and predicted (from ) values of for

Ic Ic four specimen thicknesses of aluminum alloy 2219-T87

at 76 K.

MEASURED . PREDICTED

B(cm) B(in) (kJ/m^) (^^^^) J. (kJ/m^) J" (^^^^)

Ic ic . ^ Ic Ic . z m in

.38 .15 10.3 59 13.6 77.5

.71 .28 10.0 57 ^ 13.6 77.5

1.42 .56 9.6 55 13.6 77.5

3.81 1.50 9.5 :, , 54, 13.6 77.5

Table 8. Measured values of crack extension and secant deviation

for two specimen thicknesses of aluminum alloy 2219-T87

at 76 K.

_3 Sample # B(cm) B(in) Aa(cm) Aa(10 in) secant deviation

(per cent)

C-3 1.42 .56 .002 .7 0

C-5 ' 1.42 .56 .013 5.1 0

C-1 1.42 .56 .020 8.0 2.1

C-2 1.42 .56 .076 30.0 7.1

C-7 1.42 .56 .157 62 14.0

D-7 3.81 1.5 .010 ^ 3.8 3.0

D-6 3.81 1.5 .067 26.2 4.0

D-5 3.81 1.5 ,072 28.5 5.7

132
Table Al. Jlaximum stress intensity and maximum stress intensity

corrected for the plastic zone for four specimen

thicknesses of aluminum alloy 2219-T87 at 76 K.

Sample # B(cm) B(in) K^^H^ a • ^ ) Kj_j(ksi- in"^^^) K^^* (MPa •m"'"^^) K^* (ksi- in"'"'^^)

A-4 .38 .15 46.2 42.0 57.6 52.4

B-4 .71 .28 38.1 34.7 42.3 30.5

C-3 1.42 .56 34.8 31.7 37.9 34.5

D-7 3.81 1.5 34.1 31.0 35.6 32.4

0-1 3.81 1.5 37.6 34.2 39.3 35.8

0-2 3.81 1.5 37.8 34.4 39.6 36.0

List of Figures

Figure 1. Specimen configuration used in the present study.

Figure 2. J-la data for aluminum alloy 2219.

Figure 3. Secant deviation versus crack extension for aluminum

alloy 2219.

Figure Al. Observed maximum stress intensity versus specimen thickness.

133
Figure 1. Specimen configuration used in the present study.

1,34
CRACK EXTENSION Aa (inches)

0

50

20 40 60

"T"

80 100x10

T

45

ALUMINUM 2219-T87

76 K 250

40

35 200

30

5 -

B=.15 in A B=.28 in • B=.56 in 8=1.5 in

CNI

-150 -

<

CD

100

50

0

0 .04 .08 .12 .16 .20 .24 .28

CRACK EXTENSION Aa (cm)

Figure 2. J-Aa data for aluminum alloy 2219.

1 3 5
(uig.Ol) N0ISN3iX3 >l3Vd3

CO

o

Ln CO

CM

eo _ _

CO E E

O) o u

CM «— .0^00

.-^ ^

CO

r00

h;- Oi

CM

CM ^

S CO

CO

CO

CM

OO

<

LUO

<

C\J COo

E

Z3

cd

s-o

C

o

•r-

(/) c:

O) +j

X

o

fC

s-o

I/)

I/) S->

c:o 00

CO

fO>

-a

-Mo

<u u

CO o =o

"to "c^ o =

OJ

.ti: 1- X O LU

CO

CM

13CD

CO CM OO CO CM

i"3 'NOISN3iX3 )l3Va3

136
(3/jU!-!S>|) AilSN3iNI SSBdiS lAlfllAIIXVIAI

O O O O O O CI) o CD un CO cNj o

CD in^

CO cvi

(j/^iu BdlAl) AilSN3iNI SS3diS l/\ini/\IIXVI/\l

] 37

SEMI-ANNUAL REPORT ON MATERIALS RESEARCH

SUPPORT OF SUPERCONDUCTING MACHINERY

MAGNETOTHERMAL CONDUCTIVITY

L. L. Sparks

Cryogenics Division

Institute for Basic Standards

National Bureau of Standards

Boulder, Colorado 80302

April 1976

139
Summary: Magnetothermal Conductivity

The thermal conductivity and effect of a magnetic field on the thermal

conductivity of UNS-A9 1100 in the "0" anneal condition (aluminum 1100 annealed

at 350° C for 1 hour) have been determined in the temperature range 5 k

to 20 K. A magnetic field typically increases the electronic thermal

resistance and thus lowers the total observed thermal conductivity of a metal.

For the present aluminum specimen (RRR = 32.6) the data indicate that a

6 366 kA/m (80 kOe) magnetic field reduces the thermal conductivity by 29% at

5.5 K and by 50% at 20 K. Electrical resistivity data have also been determined and show an increase in resistivity with increasing magnetic field.

Contents: Magnetothermal Conductivity

Page

1. Introduction 141

2. Procedures 141

2.1 Apparatus 141

2.2 Material 142

3. Results 142

4. Discussion 142

5. References 143

List of Figures 144

140
MAGNETOTHERMAL CONDUCTIVITY

L

.

L. Sharks

Cryogenics Division

Institute for Basic Standards

National Bureau of Standards

Boulder, Colorado 80302

1. Introduction

The objective is to determine the effect of magnetic fields on the thermal conductivity of technically important metals. The need for this information arises from the development of rotating machinery operating at cryogenic temperatures. The existing world's literature on magnetothermal

conductivity, X (H) , is concerned almost exclusively with scientific materials,

e.g., very pure materials and single crystals. A complete bibliography of

the subject was given by Sparks and Fickett [1]

.

The materials studied in this prograiri are being used or are candidates

for use in superconducting motors and generators. Optim-urr. design of these

machines, which must operate at low temperatures while in magnetic fields,

requires a detailed knowledge of how the thermal properties of the constituent

materials are affected by a magnetic field. The broad material categories

of interest include superconductor stabilizing materials such as copper and

aluminum, and structural m.aterials such as nickel alloys, stainless steels,

and metallic composites.

2 . 1 Apparatus

The principal components of the V (H) system are shown schematically in

Figure 1. Since the detailed operation of the system has been described in

previous reports [1,2] and the details of several system modifications were discussed in [3] , only a very brief reiteration of the system and its operation is included here. The specimen chamber, as shown in Figure 1, is

evacuated (operating pressure is less than 10~4 Pa) and immersed in liquid

helium. The liquid helium provides refrigeration to the specimen via the

THERMAL LINKS (capitalized terms refer to Figure 1) . The specimen is shown

mounted with its axis parallel to the axis of the superconducting solenoid

which allows the longitudinal field effect to be determined.

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. The electrical heaters were wound, one each, on the

TEMPERATURE CONTROLLED HEAT SINK (TCHS) , the SPECI2-ffiN, and the TEMPERING

POST. The pov;er supplied to the TCHS determines the approximate temperature

of the specimen; the SPECIMEN HEATER is used to establish a temperature

gradient along the specimen; and the TEMPERING POST HEATER is used to reduce

the temperature difference between the specimen and the tempering post to

less than + 5 mK. The TEMPERING POST and TCHS heaters are automatically

controlled during all tests while a constant current is supplied to the

SPECIMEN HEATER. The temperature difference existing along the sample due

to the specimen heater current is determined at the two THERMOMETER BLOCKS.

The thermal conductivity of the specimen is then computed from a knowledge

of the specimen geometry ( area/length) , the specimen heater power (Q) , and

the mieasured temperature difference along the specimen (^T) . The

2. Procedures

141
relationship of A to these parameters is given by

A = (1)

A series of measurements of AT and Q at various fields (0 :< H <. 6366 kA/m)*

and temperatures (4 < T < 20 K) result in the data presented in this report.

2.2 Materials

One material, UNS-A9 1100 (aluminum 1100), was tested during this reporting period. The mill analysis in weight percent of the stock used for the

A (H) specimen is as follows: Cu = 0.2, Fe - 0.6, Si = 0.1, and the balance

is Al . The test specimen is in the "0" condition after being annealed for

1 hour at 350° C in a vacuum of 1.33 X 10"^ Pa and then :^urnace cooled. The

DP hardness of the annealed specimen is 28 (1 kq load) and its density is

2.818 X 10-^ kg/m^ at 294 K. The residual resistance ratio, R273K/R4K'

32.6, and the ratio of area to length is 4.84 X 10"^ m. The THERI40METER BLOCKS and the SPECIMEN HEATER are attached to the speci- men by soldering as detailed in [3]. In order to avoid the difficulties

invol\7'ed' in soldering to aluminum, thin copper strips were electroplated to

the specimen as needed after the specimen was annealed.

3 . Results

The thermal conductivity of the present specimen is shown in Figure 2 as a function of temperature with magnetic field as a parameter. The reduction in conductivity from the zero field values caused by a 6366 kA/m (80 kOe)

field is 29% at 5 K and 50% at 20 K. The estimated uncertainty of the thermal conductivity data points is + 6% at T = 4 K and + 8% at T = 20 K.

Figure 3 presents the relative change in the thermal resistance, AW/Wh=o =

(Wjj - Wjj=o)/^H=0' ^ function of magnetic field with temperature as a parameter. The electrical resistivity is shown in Figure 4. The estimated

uncertainty of the electrical resistivity data points is _+ 8%.

4. Discussion

Attempts to find corroborating data have, to a large extent, been unsuccessful both for the case where H = 0 and E ^ 0. The literature contains

many references to papers dealing with the thermal conductivity of aluminum

at zero-magnetic field [4] , and several references given in [1] deal

with thermal conductivity of aluminum in magnetic fields. With two

exceptions, however, the low temperature thermal conductivity of alumi- num data pertain to very high purity polycrystalline or single crystal

specimens. This, coupled with the fact that thermal and electrical conductivities of most good conductors are extremely dependent on trace impurities

and thermal history, makes direct comparison to the present data difficult.

Powell, et aL [5] measured the thermal and electrical conductivity of

two UNS-A9 1100 specimens; one was used in the "F" condition (as fabricated)

* The International System of Units (SI) designation for magnetic field

strength is ampere per meter, and this unit is used throughout this paper.

The more conventional unit of magnetic field strength, the oersted, is

also given. Conversion of oersteds to amperes/meter is accomplished by

multiplying oersteds by 1000/4it.

142
and the other in the "0" condition (fully annealed). His results for the

"0" specimen were anomalous in that the observed thermal conductivity was even lower than that of the "F" specimen. The RRR for his "0" specimen was around 15 when the expected value for this grade of aluminum is 32. The

behavior of the "0" specimen was attributed to high impurity levels in the

stock material. Present results can also be compared with those of de Nobel [6]

In this case the data extend down to only 15 K (a 5 K overlap with the present

data) and the material is not well characterized. de Nobel's material was intended to be used in construction of laboratory apparatus (therefore, not

of extremely high purity) ; its conductivity is approximately 5% lower than

the present data.

The temperature dependence of the relative change in thermal resistivity,

AW/Wjj=Q, as shown in Figure 3 is opposite that of the copper specimens tested

earlier [7]. This temperature dependence has been observed before [8] for

single crystal specimens at lower temperatures. Further study of the present data will be necessary to examine the possible scattering mechanisms

responsible for this behavior.

The uncertainty in the electrical resistivity data shown in Figure 4 is

high due to the short specimen length and low allowable specimen current

(due to small wires) . The scatter in the zero field data was considerably

higher than the H 7^ 0 data. These data sets (H = 0 and H 7>^ 0) were taken on different days but were otherwise the same. The H = 0 curve shown in this

figure is from Clark, et al . [9]; their specimen had the same RRR as the

present specimen within experimental uncertainty. An analysis of these

Ap/pH^O data indicate that the magnetic field effect on the resistivity is

only a very weak function of temperature.

In conclusion, the thermal conductivity of aluminum UNS-A9 110 0 in the

"0" annealed condition is reduced when a magnetic field is applied. The

magnitude of the reduction caused by a 6366 kA/m field is 29% at 5 K and 50%

at 20 K. The temperature dependence of AW/%=o is the inverse of that found

for the copper specimens measured previously.

5. References

1. L. L. Sparks and F. R. Fickett; in Materials Research in Support of Superconducting Machinery - I; R. P. Reed, A. F. Clark, E. C. van Reuth (Eds.);

Nat. Bur. Stds., Boulder, CO; March 1974; AD780596.

2. L. L. Sparks and F. R. Fickett; in Materials Research in Support of Superconducting Machinery - II; R. P. Reed, A. F. Clark, E. C. van Reuth (Eds.);

Nat. Bur. Stds., Boulder, CO; October 1974; ADA004586.

3. L. L. Sparks; in Materials Research in Support of Superconducting

Machinery - IV; R. P. Reed, A. F. Clark, E. C. van Reuth (Eds.); Nat. Bur.

Stds., Boulder, CO; October 1975; ADA019230.

4. Childs, G. E., Ericks, L. W. and Powell, R. L., Thermal conductivity of

solids at room temperature and below. National Bureau of Standards (U.S.),

Monogr. 131, 603 pages (September 1973).

5. Powell, R. L., Hall, W. J. and Roder, R. H., Low-temperature transport

properties of commercial metals and alloys. II. Aluminums, J. Appl

.

Phys. 3]^, No. 3, p. 496 (March 1960).

6. de Nobel, J., Heat conductivity of steels and a few other metals at low

temperatures, Physica 17_, No. 5, p. 551 (May 1951).

7. L. L. Sparks; in Materials Research in Support of Superconducting Machinery

III; R. P. Reed, A. F. Clark, E. C. van Reuth (Eds.); Nat. Bur. Stds.,

Boulder, CO; April 1975; ADA012365.

143
8. Amundsen, T. and Sovik, R. P., Measurements of the thermal magnetoresistance of aluminum, J. Low Temp. Phys. 2_, No. 1, 197G.

9. Clark, A. F., Childs, G. E. and Wallace, G. H., Electrical resistivity

of some engineering alloys at low temperatures. Cryogenics 10^, No. 4,

p. 295 (August 1970)

List of Figures

Figure 1. Magnetothermal conductivity probe and magnet.

Figure 2. Thermal conductivity of UNS-A9 1100 "0" as a function of temperature with magnetic field as a parameter.

Figure 3. Relative change in thermal resistance as a function of magnetic

field with temperature as a parameter . Figure 4. Electrical resistivity as a function of temperature with magnetic field as a parameter. The dashed line represents H = 0 data

from Clark, et al. [9].

144
Figure 1. Magnetotherir.al conductivity probe and magnet.

145
146
H, kA/m

P 1000 2000 3000 4000 5000 6000 7000

UNS-A9II00

10 20 30 40 50 60 70 80

H, kOe

Figure 3. Relative change in thermal resistance as a function of magnetic

field with temperature as a parameter.

147
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148
SEMI-ANNUAL REPORT ON MATERIALS RESEARCH

IN SUPPORT OF SUPERCONDUCTING MACHINERY

THERMAL CONDUCTIVITY

J. G. Hust

Cryogenics Division

Institute for Basic Standards

National Bureau of Standards

Boulder, Colorado 8 03 02

April 197 6

149
Summary: Thermal Conductivity

During the current reporting period no thermal conductivity measurements

were performed. The final measurements for this program will be performed

during the next, and final, reporting period. Measurements are scheduled for

the following materials:

Maraging steel 300 (Fe-18Ni-9Co-5Mo)

Carpenter 49 steel (Fe-49Ni)

Boron-Aluminum Composite (longitudinal)

Graphite-Epoxy Composite (longitudinal, transverse)

The final report will contain the results of these measurements and a summary

of rne thermal conductivities of materials pertinent to this program.

150
I|NBS-n4A (REV. 7-73)

U.S. Ut-rl - \J ~ \^\jrnrn . BIBLIOGRAPHIC DATA

SHEET

1. PUBLICATION OR REPORT NO. NBSIR 76-839

2. Gov't Accession

No.

3. Recipient's Accession No.

4. TITLE AND SUBTITLE

t Semi-Annual Report on Materials Research in Support of Super-

' conducting Machinery

5. Publication Date April 1976

6. Performing Organization (-ode 275.03

7. AUTHOR(si R. P. Reed, J. G. Must, M. B. Kasen, H. M. Ledbetter,

!d. T. Read, R. E. Schramm, L. L. Sparks, and R. L. Tobler

8. Performing Organ. Report No.

9. PERFORMING ORGANIZATION NAME AND ADDRESS

NATIONAL BUREAU OF STANDARDS

DEPARTMENT OF COMMERCE

WASHINGTON, D.C. 20234

10. Project/Task/Work Unit No. 2756530

11. Contract/Grant No. ARPA Order 2569

; ; 12. Sponsoring Organization Name and Complete Address (Street, City, State, ZIP)

Advanced Research Projects Agency

1400 Wilson Boulevard

Arlington, VA 22209

13. Type of Report & Period Covered

14. Sponsoring Agency Code

4D10

15. SUPPLEMENTARY NOTES

16. ABSTRACT (A 200-word or less factual summary of most significant information. If document includes a significant

bibliography or literature survey, mention it here.)

Results are reported of a six-month study, ending March 1976, on candidate materials for superconducting machinery. The results cover five areas--advanced composites,

elastic properties, fatigue resistance and fracture toughness, magnetothermal conductivity, and thermal conductivity. Material properties were studied over the temperature

range 4 to 300 K. Materials studied include: aluminum alloys 1100, 2014, 2219; a nickel -chromium-iron alloy; iron-47.5 nickel; and the composite materials boron/aluminun^,

boron/epoxy, S-glass/epoxy; graphite epoxy. Some notable results of the study are:

first reports of compressive mechanical testing on composite materials at 4 K; regular

temperature behavior of the elastic constants of aluminum 2014 and 2219 and of iron-47.

nickel, which is magnetic; none of the mechanical properties of the nickel-chromiumiron alloy tested were affected deleteriously by cryogenic temperatures; in aluminum

lalloy 2219, Jjr and Kj^. are not equivalent because of sub-critical crack extension;

both electrical and thermal conductivities of aluminum alloy 1100 are reduced by magnetic fields.

|17. KEY WORDS (six to twelve entries; alphabetical order; capitalize only the first letter of the first key word unless a proper

j name; separated by semicolons

)

^Aluminum alloys; composites; elastic properties; engineering materials; fatigue; fracturjs

iron alloys; cryogenic temperatures; mechanical properties; nickel alloys; superconducti

machinery; thermal conductivity.

18. AVAILABILITY [X Unlimited

I For Official Distribution. Do Not Release to NTIS

I

' Order From Sup. of Doc, U.S. Government Printing Office Washington, D.C. 20402, SD Cat. No. C13

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