Aluminium 2. A review of deformation properties of high purity aluminium and dilute aluminium alloys

The elastic and plastic deformation behaviour of high-purity aluminium and of dilute aluminium alloys is reviewed. Reliable property data, including elastic moduli, elastic coefficients, tensile, creep, fatigue, hardness, and impact are presented. Single crystal tensile results are discussed. Rather comprehensive reference lists, containing publications of the past 20 years, are included for each of the above categories. Defect structures and mechanisms responsible for mechanical behaviour are presented. Strengthening techniques (alloys, cold work, irradiation, quenching, composites) and recovery are briefly reviewed.

Materials at l o w temperatures

Aluminium 2. A review of deformation properties of high purity aluminium and dilute aluminium alloys R. P. Reed

Pure aluminium is soft, but light, can be strengthened considerably by alloying and cold working, but is often susceptible to stress corrosion in such conditions. It has low electrical resistivity which encourages its use as a conductor. Its high strength to weight ratio, when alloyed or cold worked, promotes use in aircraft structures. It can be easily formed, thus many common items are less expensively produced using aluminium. Cryogenic technology is continuing to employ more aluminium and aluminium alloys; examples are dewars, pressure vessels, tank cars, structural supports, and normal cryogenic magnets. This wide usage reflects good weldability with retention of strength and ductility at cryogenic temperatures as well as the very significant increase in conductivity as temperature is increased. Future possible cryogenic applications of high purity aluminium, such as higher field strength normal magnets, power transmission, and stabilizing material for superconducting magnets are rapidly becoming technically feasible since recent technological breakthroughs have provided commercial availability of extremely pure aluminium (~99.99995 A1, residual resistance ratio ~ 4 5 000). For this reason it was considered timely to prepare this review which critically evaluates the mechanical properties of both high purity aluminium and very dilute aluminium alloys. Towards this end, representative data will be presented and the effects of principal variables will be discus-

The author is with the National Bureau of Standards, Boulder, Colorado 80302, USA. Received 4 April 1972. *Contrary to our normal style reference numbers throughout this paper are set in bold inside brackets.

CRYOGENICS. AUGUST 1972

sed. Considerable emphasis will be placed on data reflecting on defect mechanisms, for example, activation energies and single crystal results. As mentioned in Part 1 on resistive mechanisms in aluminium, (1)* space is limited; it is hoped that a more comprehensive review of electrical, mechanical, and perhaps thermal properties of aluminium can be published at a later date. Past mechanical property reviews have usually encompassed a vast spectrum of materials, have been of the nature of a compilation, or have been confined to only one aspect of the mechanical properties. Four prominent compilations ( 2 - 5 ) have been published in the last ten years on mechanical properties, but unfortunately these have not been updated. Furthermore, these reviews fail to reference many of the relevant papers on aluminium. One brief, but excellent cryogenic tensile property-mechanism review including some aluminium data was published by Conrad. (6) Additionally Pearson and Phillips (7) presented a review in 1957 which includes a limited mechanical property summary of all properties of high-purity aluminium. It is not intended in this review to present all data nor to present a complete reference list of past mechanical property studies on aluminium. Instead, emphasis is placed on the past 20 years; very few references to papers published prior to 1950 are listed. There are two prominent reasons for this: (1)with rare exception studies conducted before 1950 used relatively impure aluminium (usually 99.9 A1) and (2) inclusion of these earlier papers would approximately double the reference list - this was thought both inappropriate and inconvenient for this review paper.

259

Within each section property data are first presented which typify high purity aluminium. The effects of such common variables as temperature, purity, and strain rate are then documented when appropriate. Finally defect characteristics are discussed. We have chosen to emphasize the property data and the effects of major metallurgical variables on these properties as our primary intention is to supply infomaation needed in practical applications of high purity aluminium. However, the sophistication of modern-day applications increasingly demands that the user has knowledge of not only the physical and mechanical properties of materials but some knowledge of the more fundamental second-order contributions to these properties. Such contributions are clearly related to defects. Theretbre, whenever research has provided plausible information concerning defect structures or mechanisms, a discussion of defect characteristics is presented.

Elastic properties

Aluminium is a nearly isotropic metal, having an anisotropy ratio ( 2 C 4 4 / C l l - C12 ) of 1.22. The anisotropy ratio remains almost constant, varying only between 1.2 and 1.3 over the entire solid temperature range. Beginning with the early resonance work of Goens, (8) there have been a

number of investigations of the elastic coefficients. (9-19) These are summarized in Table 1 together with estimates of best values. Also included are pressure dependence and third order elastic coefficient data. Sutton, (10) Kamm and Alers, (14) Vallin et al, (I 5) and Ho and Ruoff (17)have measured the elastic coefficients over a range of temperatures (see Fig.l). The Sutton data, while showing correct temperature dependence, exhibit a large disparity in terms of absolute values. Note in Fig.1 that Ho, Ruoff data represent isothermal elastic coefficients. Included also in Fig.1 are values of the shear modulus resolved on a (111) crystallographic plane plotted in the temperature range 50 to 700 K. This modulus, G 111, can be shown to be equal to l/a (Cll - C12 + C44 ) [see, for example, Weertman (20) or Teutonico (21)]. It is less generally recognized that G 111 is independent of direction in the (111) plane. The modulus Gll 1 is important since it enters so many dislocation calculations for fcc crystals, all those that include a shear on close-packed planes. Until recently, most workers have used the bulk isotropic shear modulus, G; however, exact calculations require that the anisotropic modulus Gll I be used, and also that the temperature dependence of the cii be taken into account. Sylwestrowicz and Gibbons (22) have recommended an alternative formula which represents Gll I in terms of

Table 1. Elastic coefficients of aluminium at r o o m temperature

2nd order coefficients (adiabatic

Investigator(S)(year)

Goens

1012 dyne cm"2 TamP,K C11 C12 C44

(1933)

Lazarus (1949)

1.082

0.622

0.284

1.056

0.639

0.285

3rd order coefficients

Pressure derivatives

/aCll ~ [ \--~-/

1012 dyne cm "2 T

~,--~'/T • ~P / T C111

7.2

2.5

2.15

7.35

4.11

2.31

Cl12

C123 C144 C188

C456

Ramachandran,

1.093

0.616

0.286

Sutton (1953)

293

1.129

0.665

0.2783

Long, Smith (1954)

298

1.082

0.614

0.2848

Hearmon (1956) -review

1.09

0.63

0.28

Huntington (1958) -review

1.082

0.613

0.285

Srinivasan (1953)

Schmunk, Smith (1959)

298

1.0728

0.609

0.2832

Kamm, Alers (1964)

300

1.0678

0.6074

0.281

Vallin, et al (1964)

300

1.073

0.608

0.283

Thomas (1968)

298

1.0675

0.6041

0.2834

6.35

3.45

2.10

Ho, Ruoff (1969)

300

1.0688

0.6073

0.2832

7.22*

3.93*

2.39 *

1.07

0.607

0.283

Recommended value ~297

*Reference contains temperature dependence of pressure derivations in the

260

- 1 0 . 7 6 --3.15

0.36

--0.23

--3.40

--0.30

range77--300K

CRYOGENICS

. AUGUST

1972

O.3OxlO II

0 . 3 0 x l O t2

°26f

0.26

c

O.22 t

0.22

--

O 18 / O

10

,,?

'

'

I00

200

300 400 Temperature, K

500

? Eu

c

>. 10

1.28xlO II 1.20

°°°~%£9~^°A°°°°~A&~ [] 1 3 0 0 0

1.12(

v

O.18 700

600

1.28x1012 I. 2 0 (

E Z =

00

Cll

n

n

1.12

0

?

I.O4

E

I.O4

Z

13 n

0.96

0.96

13 [] []

0.88

0.88

o

[]

0.80

i

l

l

0.72 x IO 12

i

l

l

l

l

l

{

l

l

i

l

l

A

l

i

l

i

i

l

l

i

i

i

s-

D.BO

°oooo O0

0

0.68

0 []

0.64

u

c

l

n-n

o 17 [] E

l

O.72xi0 tl

0.6 n,

l

(

ooo~ooooooo

0.60

[]

0.64

ooo~,

& A A AA,%AA

'10

rl 2 _

? 0.60

&

0.56

E X

0.56

[] n []

0.52

0.52 [] O

0.48

I

I

I

I

I

I

I

I

I

l

I

I

I

I

I

I

I

I

I

I

I

I

A

I

I

I

I

l

I

I

I

O . 3 2 x l O I:

0.48 O . 3 2 x l o tl

° ° o ~ o c~

O.30

0.30

n 2o,

E ~'O

0.28

C44

o

0.28

[]

c >.

10

°o

O

0.26

0.26

O Sutton,99.93Al(1953)

tw

IE u

o

C44

0.24

o Kamm, Alers,99.999AI (1964} A Ho, Ruoff, 99.99AI (1969)-isothermal

¢>. 10

o

o 0.24

• Vallin, et a l , 99.995A1(1964) 0.22

% Z

n

• Best value a t 297K(se¢ Table I)

)22

[]

n

0.20

A

O

I

I

I

I

IOO

I

I

I

200

I

I

I

I

300

I

I

I

A

f

I

i 5 li 0O

400 Temperature, K

I

I

I

600

I

I

I

t

700

I

L

D.20

8OO

Fig.1 The available data on the temperature dependence of the elastic coefficients of aluminium. The upper curve represents the temperature dependence of G l l 1 [1/3 (Cll - C12 + C44)], the shear modulus resolved on the (111 ) plane. To obtain Gl11, the temperature dependent data were normalized to the average room temperature value (see Table 1 )

C R Y O G E N I C S . A U G U S T 1972

261

elastic coefficients. However, the formula suggested is only equivalent to 1/3 (Cll - C12 -I-C44 ) under purely isotropic conditions (that is, Cll - C12 = 2C44 ). Recommended room temperature engineering elastic constant values are included in Table 2. Both isothermal and adiabatic determinations are presented, along with adiabatic values calculated from the recommended elastic coefficients of Table 1. Isothermal bulk modulus values have recently been measured and compiled. (16) But, isothermalE, G, and u experimental measurements are old, dating back to Koster (31) in 1940 for E and to earlier determinations for the shear and bulk moduli and Poisson's ratio. For this reason it is suggested that the calculated adiabatic moduli, using the recommended elastic coefficient data (Table 1) and the Hill method (23) of arithmetically averaging the Voigt and Reuss moduli represent the best values. The exactness of the experimentally determined adiabatic bulk modulus, (9, 12, 17, 38, 39) Young's modulus, (39) and shear modulus, (39) and the calculated adiabatic B provide support for the validity of the recommended elastic coefficients of Table 1. Pressure dependence of E, G, and B up to 10 000 kg cm "2 (throughout the text stress and pressure units of kg mm "2 will be used. The reader is reminded that 1 kgmm "2 = 0.98 x 107 N m "2 (SI units) = 1.42 x 103 psi) have also been reported. (13, 38, 39) Wawra (24) recently presented a compilation of elastic moduli data on aluminium with references dating back to 1887. Since the latter compilation is sufficiently comprehensive forE, G, X, and v, references for these constants are omitted here. The temperature dependence of the Young's modulus, E, from several studies (25-30, 37) is plotted in Fig.2. Friedel et al (27) compare the temperature dependence of 99.96 A1 for annealed and polygonized structures. The

Table 2. Engineering elastic constants of aluminium at room temperature (recommended values in column two)

(1012 dyne c m 2 ) t

Isothermal (1012 dyne cm "2)

Expt

Calc*

Expt

0.690

0.704

0.74

0.255

0.262

0.26

0.764

0.761

0.73

Adiabatic

Young's modulus (E) Shear modulus (G) Bulk modulus (B) Shear modulus on (111) plane along any direction (Glll) Compressibility (X) =

0.249

lIB

1.32 x 10 "12 (dyne -1 cm 2)

1.32 x 10 "12 1.37 x 10 -12 (dyne "1 cm 2.) (dyne 1 cm 2)

Poisson's ratio (p)

0.353

0.349

0.34

* Calculated from 'best value' elastic constants, Table 1, using Hill

technique 1"

ldyneCm'2=l.02xl0"8 kgmm'2=l.45x10"s~i=0.1Nm "2

262

8xlO II R F r i e dete al,99.96 l At (1955) (Room temo value chosen to 25xlOI2 dYnes cm-2l 7 ecrystollized .2~,~/~ po.,,^oniz-d

8xlO

7 HowellLSf=ck!ey"

'_o6

s,okes,96o j

~s ~

tE u

4

J

"~ /~

~

'~E 4 Z

~"

/ \

~3

Be~ (1966(r)oom^tem.p...,2 \ value chosen to be 7.25xI0'" \ dyne cm-2) polycrystolline ~.

'IO

~2

J

O

5

Girard, Vidal / 995At (,960) ~

,

,

l

,

,

i

i

i

i

,

i

i

J

,

t

,

2

i

,

~)

IOO 2 0 0 300 4 0 0 500 600 700 BOO 900 IOOO Tempcroture,K

Fig.2 Temperature dependenceof the elastic modulus, E, of selected aluminium data

much greater reduction of E at higher temperatures for polygonized material (see reference 31 also) is attributed to movement of the dislocation walls aided by increased thermal fluctuations. The influences of several material or treatment variables on E (31-33) are included in Fig.3. Room temperature plastic strain of the order of 0.2% serves to lower E approximately 5%. Data on the influence of additional plastic strain (up to 85%) are conflicting; reflecting the modulus change and the increasing experimental difficulties at larger plastic strains. }COster (31) reported data for three purities (99.5 to 99.99 A1) which suggest that E decreases as the purity of aluminium increases. Recovery after plastic strain may increase E to above the annealed value, provided the strain is great enough. This is shown in Fig.3 in the variation of room temperature E with recovery temperature after 85% plastic strain. Very recently, the change of shear modulus at 4 K after neutron irradiation (34) and the change of G after neutron irradiation on warming (10 to 80 K) (35) have been measured. The results (34) presented in Fig.4 indicate that G decreases linearly with increasing defect concentration (primarily Frenkel pairs), and, correspondingly, the electrical resistivity increases until doses greater than 1018 neutrons cm "2 have been applied. The decrease in G is considerably larger than expected from calculated volume changes and expected Frenkel pair concentrations. Using the Frenkel pair resistivity of 3.4 x 10-6 ~2 cm per atomic percent (36) and assuming that the reduced resistivity is solely contributed by the production of Frenkel pairs, the reduction of shear modulus AG/Gper 1 at % of Frenkel pairs is 0.4. The recovery measurements (35) indicate that after neutron irradiation at 4 K, G initially increases (about 0.1%) in the temperature range 20 to 40 K, then decreases at higher temperatures (-0.6% at 78 K). The influence of point defects of point defect clusters is also evident from measurements of the change of E on annealing at room temperature after quenching from high temperatures. (40) Decreases (and subsequent recovery) of about 80 x 10 -6 in the value of E were observed after quenching into water from 800 K.

CRYOGENICS. AUGUST 1972

Tensile

properties

-

single

crystal

Change in resistivity, Ap, IO"9f/cm

The tensile properties of single crystals of aluminium have been examined many times since the first work of Taylor and Elam (41) in 1925. Following usual notation, several distinct deformation regions are usually detectable when the resolved shear stress is plotted against the resolved shear strain assuming [111] < 110> slip systems. The [111] < 110> slip system is customary for aluminium under normal tensile strain rates at temperatures up to about 800 K. After plastic deformation begins (defined here, for single crystals, as attainment of the critical resolved shear stress, r0), generally three deformation stages are observed in face-centred cubic metals. Related to the first two stages are distinct slopes associated with work hardening (0i, 01i) and a prescribed stress at which each stage begins (rl, rli ). The work hardening coefficient for stage one, 0i, is always less than 01i. The third and final stage is associated with dynamic recovery processes and, typically, in this stage the shear stress-shear strain slope (0111) steadily decreases. The stress, rlll, is defined as the resolved shear stress at which the slope associated with stage II, (0 II), ceases to be linear (and begins to

0

I00 "

200

300

400

I

I

I

I

500 I

600

7OO

I

I

99.999AI P3C~K =15OO 20

<~ Q30 I

40

~'I-,~L" 1-~ ~.~. 50 0

I

I

I

2

4

6

I

I

I

I

I

I

,

I

I

I

I

8 I0 12 14 16 18 20 22 24 26 28 30 Irradiation dose, lOITneutron cn~2

Fig.4 Normalized change of shear modulus, G, with irradiation dose and concomitant change in resistivity, zx p, of aluminium

t

76xlOi2

TOO.6 4.0 T

76xlO II Koster, 99.99 At (1940)

'Eu

~

7.2

r-

~-

6.8

,

, .~120 j

,

7.0t

,

A

z XPrestrain=O.15

,

,

,

,

,

,

,

~

,

,

6.8 ~"

10

1.2..4x10 H/.

7.0

5.8j 7.4x IO12

?

9.99 At(1940',

'

lq

6.2

'

' ' 5.8 0010 OOI2

7.4x10II

72

7.2 ? E

t-

Z

,%

7.0 t~-

7.0

6.B 9; .O

. . . . 9; .9 . . . . Aluminum purity)wtO/o

99199

Fig.3 Dependence of aluminium eleastic modulus, E, on (a) recovery, (b) plastic deformation, and (c) purity

CRYOGENICS. AUGUST 1972

~m

D

"1 ~ '~

I00 200 300 400

i

I

|

I

i

!

i

i

i

500 600 700 800 900

Temperature, K Fig.5 Variation of single crystal stresses TO, I"1I, and TilI as a function of temperature. Brackets indicate approximate maximum data spread from published results 4 2 - - 8 7 (1.02 kg m m "2 = 0.1 N m-2)

decrease). The stages become more distinct at lower temperatures; in aluminium, stage I is apparently difficult to detect at room temperature and above. Best published single crystal stress-strain curves include those of Hosford (42) et al at low temperatures and Garstone (43) et al at high temperatures.

Cook ¢t at 99.99 AI (IgS4)

' ~ ~ ' 0.002 O.~~ 0 4 ' 0.006 0 0'0 8 Plastic strain

0.1

-

Z

Z

'~

I I t4+L -

E

66

Hordon et at 99.6 A[(1958)

-10.2

~,,,,,

°

,

to(Scale on right)

0.8 0

6.6

6.2

1.6

5A 1000

~ --"

~

,

200 400 600 800 Recovery temperature, K Plastic strain 020 0.40 0.60 0.80 ,

",EE3.2 2.4

7.2~

'~

/

6.4 0

~.XIU

n=O.85

T

6.8

Critical resolved shear stress data have been reported by many. (42-71) High temperature (300-900 K) results (42, 43, 50) consistently indicate a gradual increase in ~'0 with decreasing temperature in the range 900 to 600 K, followed by a plateau below 600 K. Lower temperature results fail to indicate a clear trend. Fig.5 presents our interpretation of the dependence of r 0 on temperature. The large data spread results in part from material inconsistency and from differing sensitivities of the strain measurement. For example, it is commonly accepted that values obtained for r 0 or proportional limits are dependent on the sensitivity of the strain measurements; that is, higher sensitivity, such as capacitance microstrain techniques, produce lower values o f t 0 or proportional limit. It has

263

also been demonstrated (47, 60, 63) that r 0 is higher for less pure specimens. Although data are scarce, it was thought instructive to present these in Fig.6. As also illustrated by hardness data (presented in a later section), the strength dependence becomes less sensitive as the purity increases. Zone refining of 99.992 AI also produces no detectable change o f r 0. (71) Substructure has a decided influence on r 0 values; r 0 increases with decreasing subgrain size. Subgrains can be easily introduced into aluminium by straining at room temperature between 5 and 20% followed by annealing at rather low temperatures (400-600 K). The influence of substructure on r 0 is discussed and documented in the last section on strengthening. Additionally, there are limited data (59) which imply that r 0 increases in proportion to the density of growth or forest dislocations, as determined from etch pit density counts in 99.99 AI. The influence of other strengthening agents, such as dilute alloying additions, (44, 4 9 - 5 1 ) is considered in a later section on strengthening. Testing under reduced pressures or under conditions in which the surface is continuously removed does not influence r 0 (56); however, ~'0 is increased if the specimen sizes are decreased (89) or if the oxide layer (formed during annealing) is not removed. (90) Irradiation has considerable influence (86); neutron irradiation of 1.5 x 1018 n cm "2 at 20 K increases r 0 from about 0.2 kg mn "2 to 12.5 kg mn "2 ! Estimates of the stresses, I"ii and Till, are plotted as a function of temperature in Fig.5. Little data are available for TII and since considerable scatter exists in the data, only temperature-dependent investigations (43, 52, 74) were considered in the dubious estimation of T-iipresented in Fig.5b. Our rll I estimate from all reported values (42, 43, 46, 49, 52-54, 60, 65, 7 4 - 8 7 ) tends to follow data from crystals having an approximate <123> orientation (therefore an initial single slip system). The ril I values of crystals having <100>, <110>, or <11 I> directions parallel to the tensile axis tend to be higher at temperatures above approximately 150 K and lower at temperatures below 100 K; that is, the rll I values for these low indices directions exhibit less temperature-dependence. If the curve of estimated rll I temperature-dependence in Fig.5b is plotted using a log e o ordinate axis, two linear ranges are very apparent: (1) a region from 0 to 145 K, having the higher slope; (2) a range above 145 K, having the lower slope. Lack of comparable documentation makes a similar analysis o f r 0 and TII data probably useless. The magnitude of 7"111has been shown to be reduced when testing under reduced pressure (56) and flow stresses are reduced if the surface layer is continuously (84) or periodically (85) removed by chemical or electrolytic polishing techniques. Although care has been exercised in extracting these data from published curves and tabulations, aluminium single crystal stress-strain curves do not exhibit extensive stage I strain, they have in many cases a non-linear stage II, and stage III begins at relatively low stresses and strains. These traits undoubtably account for the considerable variation in reported data and should stir the reader to caution in placing too much validity on data synthesis from published reports. Work hardening rates of stages I and II are orientation dependent, orientations which produce slip on more than one system having higher rates. The temperature dependence of a few work hardening rates (selected on the basis of consistent available data) (42--44, 46, 47, 60, 61, 65, 72-87,

264

4.0

3.2

?

2.4

1.6

0.8

0

99.0

I

I

99.9

99.99

99.999

Aluminlurn purity

Fig.6 Dependence o f TO on specimen purity f r o m published data 47, 60, 63 (1.02 kg mm -2 = 0.1 N m-2)

4O

32

";'E

24

OK~ (i00)

16

O H ~

eI I

0

20

I

(123)

< 123 ) I

60

I

I

I00

I

I

140

I

J

180

I

=

220

I

I

I

260

300

Temperature, K Fig.7 Temperature dependence of stage I and II w o r k hardening coefficients of selected orientations. Curves are estimated f r o m available data (1.02 kg mm "2 = 0.1 N rn-2)

156) is shown in Fig.7. Contrary to what is sometimes pro-

claimed, 01i is not only orientation dependent, but is also quite temperature dependent (particularly below approximately 100 K) and pressure dependent (reduced by a factor of about two when the pressure is lowered to about 5 x 10-8 mm Hg). (56, 139) The stage I work hardening rate is also slightly reduced when crystals are tested in atmospheres of decreased pressure. (56) Although only one data point is available at 23 K, (73) it appears that 01 is only very slightly temperature dependent. Both 01 and 0ii are dependent on surface conditions; surface removal during testing reduces 01 and 011 by as much as one-half. (85, 89) Slip or glide processes may be monitored by microscopic observations of polished surfaces of single crystals as crystal deformation proceeds. During stages I and II in crystals oriented for a single glide system (inside the unit stereographic triangle, preferably in the vicinity of < I 23>) slip lines on one set of [111] planes are predominantly observed. The appearance of cross-slip (slight offsets from slip line, may be associated with wavy slip) has been associated with the onset of stage III (Till). (49, 55,

CRYOGENICS . A U G U S T 1972

80, 82) The onset of cross-slip has been found to correspond to about 0.5 kg mm "z for 99.99 A1 crystals at 196 K (80) and to about 0.4 kg mm "2 for 99.992 AI crystals at room temperature. (55) These fall well within the data spreads for these temperatures shown in Fig.7. The stress-strain characteristics of aluminium crystals with orientations parallel to the tensile axis such that more than one slip system is operative have been examined (42, 74, 77) and discussed. (87, 88) It is pointed out that two distinct multiple slip orientations are evident, < 1 1 1 > and <100>. Although the <111> has 6 slip systems, compared to 8 for <100>, it exhibits flow stresses nearly twice as large as the latter orientation. Polycrystalline stressstrain data generally fall between these polyslip extremes. Tensile

-

Polycrystalline aluminium tensile properties have been measured over a period of many years and some have been previously compiled. ( 1 - 6 ) Behaviour of polycrystalline aluminium under applied tensile stress differs distinctly from the behaviour of single crystals of the same purity. For example, easy glide, characteristic of stage I deformation, is usually absent in polycrystals. However, when specimens contain less than approximately eight grains per cross-section, easy glide may sometimes be observed at very low temperatures (that is, 4 K). The addition of grain boundaries serves to strengthen. This strengthening role of grain boundaries is discussed in greater detail later in this section. Characterization of polycrystalline stress-stress curves usually involves the documentation of (1) the stress at which the strain dependence on applied stress ceases to be linear (plastic or proportional limit), (2) flow stresses (OF) at prescribed plastic strains (conventional yield strength is taken as the stress at 0.2% plastic strain), (3) the highest stress which the specimen withstands at a given strain rate (the ultimate tensile strength or, merely, tensile strength), (4) the total elongation which the specimen undergoes prior to fracture, and (5) the reduction of area at fracture, that is, the specimen cross-sectional area at fracture compared to the original cross-sectional area. In aluminium these characteristics are influenced considerably by temperature, rate of specimen strain, grain size, purity, and dislocation density. Considerable variation in reported tensile data on polycrystalline aluminium is apparem when the literature is appraised. There are several reasons f-~ these variations. In some cases metallurgical control of grain size has been neglected. In others, particularly in the 99 to 99.9 purity range there is evidence that impurity distributions, controlled by thermal history, were not random and thus influenced the tensile properties. But one dominant uncontrolled variable in many studies appears to be the dislocation density. Dislocations may be inadvertently added to aluminium specimens very easily by the following methods: (1) specimens are easily bent, particularly prior to the initiation of low temperature tests, (2) rapid cooling results in excess vacancies which, under sufficiently short times, may coalesce into platelets which subsequently collapse into dislocation loops, and (3) differing solidification techniques and purities result in variations of intial dislocation densities. Another sensitive variable is the specimen surface. All thermal treatments invariably result in an oxide layer on

. AUGUST

Temperature

dependence

It is commonly thought [for example, see Conrad (S)] that the flow stress, r, is composed of two components; r* being dependent on temperature and strain rate while r# is temperature independent when corrected for the dependence of the shear modulus on temperature, hence r = r* + rkt

polycrystal

CRYOGENICS

the surface. Flow stress values are directly proportional to the depth of this oxide layer. To achieve reproducible results it is necessary to reduce this oxide layer to a consistent minimum by chemical or electrolytic polishing immediately prior to testing. Careful, consistent control of the surface is essentialto ensure reproducible tensile data for aluminium.

1972

(1)

It is thought that the process of low temperature glide is thermally-controlled with fluctuations of thermal energy supplying the necessary energy increment to assist dislocations past obstacles that have short range stress fields. Examples of such obstacles are forest dislocations (those not taking part in the slip process), impurity atoms, and numerous dislocation reactions; these require only small activation areas and volumes. Therefore, the term, r*, is critically dependent on temperature. The temperature independent or thermal flow stress component is thought to represent the stress necessary to overcome obstacles that have larger, longer range, stress fields. These obstacles such as precipitates or dislocations on other slip planes require activation energies too large to be supplied by thermal energy and thus are relatively insensitive to temperature variations. The reversible temperature dependence, that is, the temperature dependence of a material having the same defect structure, has been measured.(52, 9 1 - 9 7 , 107) Measurements are normally conducted by stressing to a prescribed plastic strain, then suddenly changing the temperature. The results from these tests, plotted in Fig.8, are remarkably consistent and portray three distinct temperature ranges; a steep temperature dependence below 140 K, arange between 140 K and about 450 K that is essentially independent

1.5 1.3

A Hirsch,Worrmgton(19611 _

~

I.I

C

"trel. S "

119551Basinski119591

G BuIlen(101")51 H Boll11957) E G

09

A ~

0.7 0.5 O

~ i IOO 2 0 0

~ 300

I 400

= 500

= 6OO

= 700

= 800

900

Temperotunt, K Fig.8 Reversible temperature dependence of flow stress corrected for the temperature dependence of the shear modulus, where "rT is the shear stress of temperature T, 1"300K is the shear stress at 300 K, G300 Kis the shear modulus at 300 K and GT is the

shear modulus at temperature 7". Cottrell and Stokes,91 Hirsch and Warrington,92 Basinski,93 amd Mitra, et al94 results contain single crystal measurements

265

of temperature, and a high temperature range of decreasing flow stress with increasing temperature. The high temperature range is thought to be associated with diffusion controlled processes as all measured activation energies (see later section) are in the neighbourhood of the activation energy for self-diffusion. The range in which the flow stress is nearly independent of temperature represents the athermal component (zu) of the flow stress. At lower temperatures thermal energy is needed to assist glide dislocation motion past local defect barriers. Results of these experiments (52, 9 1 - 9 7 , 1 0 7 ) show to a first approximation, the ratio of flow stresses at two temperatures (determined by rapid temperature change at constant strain) is a constant, independent of the amount of plastic deformation. This is commonly referred to as the CottrellStokes law. (91) Recently, careful stress measurements on aluminium single crystals (179)have indicated that this law is not exactly correct; extrapolation of fT. 7"T versus r :~. 1 (between about 77 and 82 K) reveal a positive non-zero intercept. Previous aluminium temperature cycling measurements (180, 181)between 293 and 77 K were interpreted as producing a zero intercept.

25 20- ~ 15 IO

//Ultimote tensile strength, 99.0- 99.9AI

L_ ~

Y/~ldstrength, 99.O-99.9At Yield strength

/ ~

S ~ ~ . 9 9 - 9 9 - 9 9 9 A 1 I

O

I

IOO 200

I

I

/

~

I

300 400 500 600 700 800 Tempcrflture, K

8O

-

On the other hand, a large number of investigators have determined absolute values for yield strength, tensile strength, etc of aluminium. (98-184) These absolute data have a great amount of spread, for example, the variation of yield strength (0.2% offset) at room temperature varies from 0.35 to 10 kg mm "2 . Therefore, at this time it is very difficult to provide reliable estimates of appropriate values. In Fig.9 guesses of the temperature dependence of the yield strength (0.2% offset) for two classes of purities and of the tensile strength for commercial purity aluminium are presented. Undoubtably, the experimental problems of preventing the addition of dislocations (through specimen bending and misalignment, etc), of dissimilar surface conditions, and lack of solute segregation control have enormous effects on the early flow behaviour of high purity aluminium. Reproducibility of the relative reversible temperature-dependent flow stresses is significantly easier since specimen prestraining tends to diminish the influence of these variables and to establish a uniform defect structure. Notice in Fig.9 that the temperature dependence of the tensile strength is considerably greater than that of the yield strength at low temperatures. An inconsistency appears in that the low temperature yield strength data do not indicate a greater temperature dependence, which should follow from the greater single crystal r 0 and rll I dependences. Therefore, it is entirely possible that the low temperature yield strength data in Fig.9 are underestimated.

Effect of grain size. Although there have been a few deformation studies on aluminium bicrystal specimens, ( 1 5 1 155) the most illuminating information pertinent to the role of grain boundaries in strengthening have come from studies in which the grain size was varied. (87,101,109, 114, 122, 156-158) There have been recently several excellent reviews, theories, and suggestions concerning the role of grains and grain boundaries in polycrystalline deformation. (158-160) Most data relating flow stress to grain size at constant temperature appear to follow the Hall.Petch relationship o f = °o + k d - ~

266

(2)

40 20[O

,

,

,

Tensile elongation :o frac:ure, 99.O- 99.9 AI.

IOO 200 300 400 500 6OO 700 800 Temperature, K

Fig.9 The variation with temperature of polycrystalline tensile yield strength and percent elongation (tensile). Curves are typical and are thought representative of published information86-181 (1.02 kg mm'2 = 0.1 N m-2)

where d is the average grain diameter, and k and o 0 are material constants depending on purity, temperature, and strain. However, in Fig.10 the log e d is plotted against of using available data. (101,109, 114, 122, 176,177) Perusal of data points from Carreker and Hibbard (101) suggests that oJf~ kd "'/2, but the data from the more recent studies (114, 1"22, 176, 177) appear to be best-fitted by other functions. The reason for the lower dependence on d of the Hultgren data (114) is not obvious. All experiments used material initially heavily deformed at room temperature and subsequently obtained their grain size variations by recrystallization at a series of times and temperatures between about 570 and 860 K. The primary variable between the studies appears to be the purity, Hultgren used 99.9 A1, while 99.99 A1 was used in the other two studies. Also, Hultgren was careful to identify subgrain contributions (see later discussion on Substructure strengthening), while the others have not documented variations of subgrain and structure in their specimens. The influence o f d on t~f at lower temperatures is inconclusive from available results. (101, 114) Knoll and Macherauch (122) indicate a slight increased dependence of of on d at 76 K compared to 295 K while the Carreker and-l-Iibbard data portray a decided decrease in dependence of of on d at 20 K. Contrasted to that found in some investigations on other materials, the Carreker, Hibbard results indicate a greater sensitivity at lower strains. It has been argued for some time now that either grain boundaries act as barriers to dislocation motion, producing pile-ups and thus hardening, or that grain boundaries act as nucleation sources for dislocations, thus increasing the rate of work hardening by dislocaton addition. The pile-up approach results in (2) where k is a function of the stress

CRYOGENICS . AUGUST 1972

necessary to nucleate dislocations in the grain adjacent to a pile-up and is a function of the shear modulus, the Burger's vector, and an orientation factor. Li and Chou (161) have considerably extended these basic considerations. The dislocation density model assumes that the expression

9 8

?E

"

9xlO-7

•

P ~ ] r s o n , Phillips [1957) / Gok - t

. Moreou, Mudrovo,

6

6 (~1 'E

E

of= o 0 + otGb p a/2

(3)

has experimental validity. In (3) p is the dislocation density, o 0 and a respectively represent a constant lattice frictional stress and a numerical constant ranging between 0.2 and 0.8, G is the shear modulus, and b is the Burger's vector. In materials other than aluminium it has been shown that p ~ 1/cl; using this proportionality, (3) becomes functionally equivalent to (2). Modifications to these two basic approaches are currently being proposed. (159-161) Another role of grain boundaries is to provide a sink for impurity atoms. By proper thermal treatment a limited purification of the solid solution maxtrix may be obtained by segregation of impurities to grain boundaries. It is thought that by changing the relative amounts of impurities in solid solution or at grain boundaries that secondorder changes of mechanical strengths can be achieved. To date no data on the effect have been published. Data on the increase of fracture stress at 4 K with decreasing grain size are available. (157) Closely connected with the effect of grain size on strength is the influence of substructure size and orientation, these are discussed later in the section on Substructure strengthening. Effect o f purity. Judging from 7"0 dependence on purity (Fig.6), one would expect that the flow stress (of) would decrease with decreasing impurity content. Data on the influence of random impurities on of are presented in Fig. 11. (7, 69, 1 0 6 , 1 1 1 - 1 1 3 , 1 2 7 , 1 3 1 , 1 5 0 , 177) Although no consistent slope is apparent (this is understandable in

3,2 ~

=O.OO2, Hansen [ 1969]

2.8 2,4

~

k

/ ' = 0 " 0 7 0 / C a r r e k e r Hibbord l,gsl]

'

? 2.0 •~

Yr.6

~

~

\

~ ~

/ , ¢=0.01 l Jooul. .. ¢ =O.OOSJBricot

1~1.2

=o.o2s, 08 O,4

-

Hultgren Cl96~,}

\ \ \~\ \ ~k

~ , Knol I, / ¢ =0.03~ Macherauch

\\

1,9691

/Elastic ~ M a c h e r a u c h

O O.OI

I

I

I I I IIII

I

I il,liil

O. I

I

I.O

I

limit

I I

[ 1961] III

IO

Average grain diameter.ram Fig.10 The dependence of flow stress on average grain size at 298K (1.02 kg mm-2 = 0.1 N m "2)

CRYOGENICS . AUGUST 1972

~

5

g

4

if)

3

Howe, fl9531 - - / ~,~

~ Nook (1966)

0

RIvslnllg~l ) L

99

~

o '(\

~

(f960)

/

I

z

4

7

Correk¢~

w'~,~.~.~l~H¢llin9

99.9

5

\

.0

I

99.99

Hibbord (1957)

' Ncunzig (1957)

~:..~v

I

I

99.999

99.9999

O

Aluminium purity Fig.11 The dependence of f l o w stress (solid lines) and tensile stress (dashed lines) on aluminium purity at 298 K. Flow stress data are at plastic strains of 0.001 t o 0.01 except for the microstrain measurements (e = 1 x 10 "6) of Rosenfield and Averbach

view of the different specific impurity types which are probably present in specimens of separate investigations), it is clear that higher purities have lower flow stresses. The 10-6 e offset data of Rosenfield and Averback (69) imply that flow stresses in the microstrain region are not as sensitive to purity as are of values at larger strains, after considerable work hardening has occurred. Work hardening dependence on purity has been experimentally identified (178) with data presented which imply that larger dislocation densities generated in the less pure specimens are responsible for the observed dependence. Available date (7,173) indicate that there is a reduction in tensile strength of higher purities, stronger than the dependence of yield strength on purity. Maraev, et al (173) have shown a linear dependence of tensile strength on the log of specimen residual electrical resistivity. This correlation between electronic purity (a measurement device to obtain estimates of total solute impurity) and flow stress is presently being studied in our laboratory. It would be useful to know the influence of purity on the temperature dependence of the flow stress, that is, whether 7"* or 7"/.tis dependent on purity. There are limited data (178) which indicate that 7"* is dependent, increasing with increasing impurities. If purity does significantly affect 7-*, then the slope of the reversible flow stress curve at low temperatures (Fig.8) would be altered as a function of purity. Unfortunately all data have been obtained on specimens of similar purity. Determinations of the effect of purity on single crystal parameters (other than 7"0) and activation factors are almost non-existent. Strain rate dependence Several (79,164, 165) have recorded stress-strain behaviour of aluminium single crystals at high tensile strain rates (103 s"l) (164) or under tensile-impact conditions. (79, 165,169) Apparently r 0 is increased, the strain range of stage I is lengthened, and 0ii and 7"111are increased by an increase in ~. Dislocation cell structure, per unit strain, is found to decrease at higher strain rates. (168)

267

Fortunately (or unfortunately for the reviewer) commercial purity aluminium has been chosen as a primary test material in a number of high strain rate (up to 120 000 s"I) test programmes (98, 104, 110, 111,128, 166-175, 182, 184) even though the selection of aluminium for such studies has been criticized on the grounds that aluminium is not as ratesensitive as many metals. Strain rate studies of aluminium have two thrusts: (1) to determine activation areas which may provide insight into the dislocation mechanisms involved and (2) to provide engineering data, usually at temperatures above room temperature, as required for specific applications. The activation area studies are not specifically discussed in this review. Two types of functions relating stress (of.) to an imposed strain rate (~) have been considered for ~luminium. Early data (171,172) were thought to be best represented by the power law

0.4

0.3

D~c~i~fstcr,MoIvern [ 1963] Hockctt[ 1967] Aider.Phillips [ 1:__50~005?

I

/

•~ . ~ O. 2

x

/-"¢=0020

0.1

o" =

¢,~

X

o.f o 0 ~n =

(4)

with n regarded as a rate-sensitivity parameter. Values for n and o 0 have been presented (172) for the temperature range 295 to 573 K for ~ values between 0.4 to 311 s"1 . Later results (104, 111) suggest that a better fit is obtained using the logarithmic dependence

of = o 0 + klog~

(5)

where the fraction k/o 0 may be related to rate-sensitivity with o 0 representing of at strain as illustrated in Fig.12 which summarizes the-data of several investigations (98, 104, 111) using ~ ranges from 0.1 to 2 000s "1. At high temperatures the dependence of of on ~ increases rapidly. The k/o 0 data below about 400 K'are inconsistent. As the lower temperature data were obtained using the same techniques by the same investigator, this implies that the logarithmic function is only appropriate for use at higher temperatures, where, incidentally, diffusion is thought to be rate-controlling, Another point to be taken from Fig.12 is that at higher strains at the same temperature the ratesensitivity parameter, k/oo, is lower and its dependence on temperature is less.

Other variables Lowering the pressure has been shown to reduce the rate of work hardening. Testing at 8 x 10-7 mm Hg at 295 K has been found to reduce the flow stress by about 30%, (138) testing under a hydrostatic pressure of 12 000 kg cm "2 increased the flow by about 20%. (185) Presumably this effect is related to a decrease in specimen surface stress under lower pressures. Increased flow stresses have been obtained for smaller diameter specimens, (125) other parameters being identical. This specimen size effect has also been accounted for by variation in the relative contribution of the surface layer stress. Irradiation has been demonstrated to considerably increase the yield strength. A fast neutron dose of 7.6 x 1020 nvt, sufficient to produce cavities in aluminium, results in a 250% yield strength increase. Additionally, astonishing increases in low temperature tensile strength have been obtained after neutron irradiation. For example, at 27 K the tensile strength of commercial and super-purity aluminium was found (181) to increase by a factor of two with a corresponding sharp decrease in

268

0

I

200

i

i

400 600 Ternpcn0tur¢,K

i

800

I0 0 0

Fig. 12 The variations of parameter k / o 0 o f the stress--strain rate function, o f = o 0 + k log ~, as a function o f temperature. See text f o r further clarification and discussion

elongation (from about 30% to 2%) after exposure to a neutron flux of about 1019 n cm "2.

Fatigue When a solid is subjected to alternating or reversed stresses (or strains), instead of continuously-applied tensile, compressive, or shear stress, the peak cyclic stress which causes failure is considerably lower than the tensile, compressive, or shear ultimate strength. Under fatigue loading conditions the number of cycles needed to promote fracture (fatigue life) becomes less dependent on stress amplitude at lower cyclic stresses. In some cases the fatigue life at low stresses levels has been found to be essentially independent of stress or strain amplitude (typically plotted as S - N curves); in Fig.13 the early fatigue data of Moore (186) and McCammon and Rosenberg (187) illustrate this. Although other investigations (29,188-198) do not indicate the existence of a clear endurance limit, an approximate room temperature endurance limit (arbitrarily taken as the fatigue strength at 108 cycles) for 99.95 A1 is 1.0 kg mm "2. Under tension-compression loading conditions it is clear that aluminium fatigue life is extremely dependent on temperature and mildly dependent on grain size and rate of cycling. Although no data are available to suport this contention, it is probable that the fatigue life is also mildly dependent on purity, increasing with increasing impurity concentration. Fig. 13 contains selected fatigue data (29, 187-192, 250) for annealed aluminium. Included are the very unconventional measurements of Borovik, et al (254, 255) who wished to improve the ability to predict stability and lifetime of aluminium-wound cryogenic magnets. Their data were obtained by merely pulsing a small aluminium wire magnet, calculating the tensile stress from current density, field strength, and wire radius. Reports of stress (strain) versus cycles to failure not included in Fig.13 include the early work of Moore, (186) who indicates an endurance

CRYOGENICS. AUGUST 1972

limit of 7.3 kg mm "2 at room temperature, the data of Greenall and Gohn (193) which show an endurance limit of about 6.6 kg mm °2 for cold worked 99.6 AI, the stresscorrosion (salt water) results of McAdam, (194) the Hordon and Wright (195-197,253) data on the effect of partial oxygen pressures and temperature, the results of Zambrow and Fontana (149) on hardened 99.6 AI, the constant strain (2%) fatigue data from 298 to 787 K of Blucher and Grant, (198) and the room temperature continuous and randomly applied peak stress results of Hod et al. (247)

3"0 x108 * .

~

4.5 2OK

4.o 3.5

3.0 (41

2.5 2.0

sl

4

~

90K

I.O ~1 ~ L

t Ill

03

[m °aoO

Io'

Io 2

~3

io 4

523K

-b

12a) jo s

io6

io7

io 8

Number of cycles Fig.13 A l u m i n i u m strength versus number of cycles to failure, Curve (1) - Awatini et al (1969); 99.999 AI, 17.7 kHz; Curve (2) - Mauer, Weiss (1969), 99.98 AI, (a) 160 Hz, (b) 20 Hz; Curve (3) - Daniels, Dorn (1957), 99.998 AI, tens - comp 24 Hz, (a) 1.0 mm gs, (b) 0.1 m m gs; Curve (4) - McCammon, Rosenberg (1957), 99.99 AI, tens -- comp -- 225 Hz; Curve (5) Halford, M o r r o w (1962), 99.6 AI, torsion; Curve (6) -- Thompson, Backofen (1971), 99.6 AI, tens -- comp -- 30 Hz, (a) 2.0 mm gs, (b) 0.2 mm gs; Curve (7) - Howell, Stickley (1952), 99.6 AI, tens -- comp; Curve (8) -- Zamrik, Shewchuk (1967), 99.993 AI, 30 Hz; Curve (9) - Borovik, et al (1965, 1966), magnet wire pulses, tens (1.02 kg mm -2 = 0.1 N m "2)

CRYOGEN ICS . A U G U S T 1972

,

I00

,

,

200

,

~, " , 300

,

400

,

,2

5OO

106). The data are from McCammon and Rosenberg. (187) Zambrow and Fontana, (149) and Daniels and Dora. (190) It is apparent that the temperature dependence of the fatigue life changes abruptly at about 150 K. The sudden steeper temperature dependence below 150 K corresponds remarkable well to the temperature dependence characteristics of both the flow stress (Fig.8) and 7"ii1 (Fig.5). This suggests that the mechanisms responsible for the flow stress-temperature dependence also play a dominant role in fatigue fracture. It is not possible to predict temperature-dependence trends at higher temperatures from the meagre data available above 300 K.

5.0

(

,

Fig.14 Fatigue strength versus temperature for aluminium (1.02 kg mm "2 = 0.1 N m "2)

5.5

~

]

Zombrow,F - o n = c°ldwOrked 99'6 At

Tempercltunz, K

6.0 x I08

=

I

O

The fatigue strength is plotted in Fig.14 as a function of temperature at constant number of cycles to failure (105 ,

E

IOScycles ~lQ'cyclls

I [

s

Major variables

'E

6 R) cyck~ McCommon, Rosenburg (1957) 99.99 A[

2.0 ~•cycies L ~.~

Fatigue studies are extremely numerous, dating from the late 19th century experiments on the 'patience' of railroad rails. The rather complete reviews by Kennedy (199) on metallurgical parameters and by Forrest (233) on properties and applied problems nicely summarize many of these studies. Earlier general reviews include those of Glover et al, (231) Cazaud, (232) Thompson, (237) Gough, (239) and Benham. (249) Several of the more recent conference reports on fatigue include those of Sines and Waisman, (234) Parker and Fegredo, (235) and the 1955 Stockholm colloquium, (240) the Orlando, Florida conference, (241) and Mott's 1957 discussion of work hardening and fatigue. (242)

o~

~

io

A significant improvement in fatigue understanding has been the correlation of atmospheric effects with fatigue crack propagation and oxygen adsorption at the crack tip. It has been known for some time [see original review of Gough, Sopwith (200)] that increased resistance to fatigue fracture could be achieved by testing in reduced air pressures. Additional confirmation that fatigue life in aluminium could be increased by a factor of 4 to 15 was provided by Wadsworth, (201) Kramer et al, (124) Shen, (252) and Ham and Reichenback. (256) Wright and Hordon (195-197,253) measured crack propagation rates at two temperatures and cycles to failure at two frequencies as a function of air, oxygen, water vapour, and hydrogen pressure. It was found that reduced pressures of all environments improved fatigue life with a maximum life of about 2.5 x 106 cycles for commercial 99.6 A1 grade achieved at reduced pressures and 25 or 50 Hz. (253) They concluded that environment has little effect on crack nucleation but that fatigue crack propagation is significantly influenced by two parameters: (1) the rate of oxygen adsorption on the crack (proportional to Po/T, where P0 is the oxygen partial pressure and T is the temperature) and (2) the rate of crack propagation (proportional to T for constant applied strain increment). This second criteria, the dependence on crack propagation rate, was also confirmed by studies on oriented single crystals. (236) If cracks propagate too fast, oxygen adsorption requires an increased oxygen partial pressure. (195-199) A good review of gaseous environmental effects on fatigue characteristics has recently been written by Bohmer and Munz. (251) Consideration of corrosion-fatigue should help ghed light on studies relating specimen surface, specimen geometry, or test conditions to fatigue life. Tests on aluminium plated with copper and silver (202) revealed that copper-plating reduced the fatigue life of aluminium while silver plating resulted in increased fatigue life. This may reflect the increased susceptibility of copper to atmospheric corrosionfatigue while silver would be expected to be less suscept. ible, (201) both compared to aluminium. The longer

269

I

fatigue life from higher frequency tests of frequencydependent tests (191,257) (see Fig.13) also may be attributed to oxygen adsorption on cracks. Even though the vibrational, near-ultrasonic tests require more cycles to failure after crack initiation, the elapsed time is shorter (and the crack propagation rate greater) when compared to typical 20--200 Hz tests.

¢

Failures in aluminium under fatigue loading conditions involve an initial rapid strain hardening, deformation (crack initiation), crack growth, and, finally, crack propagation. To attempt to characterize these processes, a number of theories have been proposed and a large number of optical and electron microscopy and x-ray diffraction studies have been reported. Most of these have been reviewed either by Thurston, (204) Low, (205) Kennedy, (199) or Zackay et al. (206) Later papers (207-230) also have contributed to the understanding of aluminium fatigue failure or fatigue defect structure. Very briefly, for alunrinium the following characteristics have been identified. As a result of constant, reversed strain cycles, dislocations begin clustering almost immediately along [111], [100], and [110] planes. This clustering produces dense dislocation tangles aligned parallel to the active slip planes. (227) Such tangles gradually evolve into a cell structure, the rate of evolution depending on whether double or single slip within each grain is responsible for the strain. (209) This cell structure tends to be elongated underneath surface extrusions and intrusions. Considerable evidence indicates that the cell formation process is greatly enhanced near the specimen surface. Cell volumes gradually diminish with increasing strain to a limiting size. This limiting subgrain size is dependent on strain amplitude (209): high strain amplitudes producing the smallest cell size with smaller.amplitudes resulting in a much larger subgrain size (see Fig.15a). Typical average dislocation densities are about 109 lines cm "2. On continued cycling this density increases very little. Two types of dislocation groups are observed: (1) cell walls, composed of tangled and jogged

270

-

0

.

0

0

3

,

¢=0.002, 298 K ~

2 I

0101

I

Io°

I

I

lo'

I

~ lo3 Cycles to failure

I

Io 4

I

Io s

,o6

Sol, rain volume. 1(59cmz ,o'

i01

I

,o

,o

I

I

, oo

c

j~J

u

99.99 At

~

t~ 4

Defect processes

¢ -o.002, 77K

3

During the fatigue process extensive localized upheavals (extrusions) and identations (intrusions) gradually develop on the surface. Correlations with internal defect structure have recently been made (see later discussion). Surface damage always has been observed to preceed internal damage. In fact, polishing after partial fatigue to remove the induced damage in the surface layer will either restore the fatigue life (201) or lead to an extension of fatigue life in low cycle fatigue. Fatigue life is also sensitive to grain size; at low temperatures smaller grain sizes result in greater fatigue resistance. For aluminium this dependence is shown in Fig. 13. Further. more, Ray and Seal, (203) using a number of different grain sizes (obtained through combinations of room temperature plastic strain and 673-773 K annealing temperatures) found that the log of the number of cycles to failure was directly proportional to the log of the reciprocal of the average grain diameter. Presumably grain boundaries act as obstacles to crack propagation and ensure more diverse crack paths, which must be followed through a finer grain structure. At higher temperatures it has been shown (190) that the fatigue strength is insensitive to grain size variations.

Alden Backofen(1961), 99.997 AL

4

I

0 t

,o°

I

I

lo'

lo2

I

I

,o

io'

I

,os

Cycles to failure Fig.15 a -- Energy needed t o d e f l e c t a constant strain i n c r e m e n t versus cycles t o f a i l u r e f o r a l u m i n i u m ; b -- A m o u n t o f plastic d e f o r m a t i o n per fatigue cycle versus cycles to failure ( C o f f i n 225) or subgrain size (Grosskreutz, W a l d o w 209) -- at r o o m t e m p e r a t u r e for a l u m i n i u m

dislocations and (2) loop patches, comp.osed of a high density of dislocation loops. The occurence of loop patches is also very dependent on strain amplitude. At low amplitudes loop patches predominate. Isolated oscillating dislocations are thought to generate the loops. Presumably in grains oriented for single slip at high strain amplitudes there exist sufficient dislocations to sweep and annihilate these loops while at low amplitudes there are not enough glide dislocations available to annihilate all of the produced loops. Loop densities of 5 x 1014 cm "3 and loop diameters of 900,8, have been estimated. However, under cyclic, but not reversed, tension at 78 and 300 K a disproportionately high density of dislocation loops (compared to static tensile) was not observed. (211,219) Additionally these'same investigators reported no detectable differences between dislocation structures produced by static tension, static creep, and cyclic (non-reversed) tension. To summarize, it is evident that reversed strain fatigue, through the action of oscillating dislocation interactions, produces a heavy concentration of point defects and a

CRYOGENICS. AUGUST 1972

tangled, jogged dislocation subgrain structure. The production of both point defects (resulting in a large loop density at 300 K) and substructure is dependent on strain amplitude, number of cycles, grain orientation, and, probably, on temperature. Fatigue cracks

Most fatigue cracks are thought to originate at or near the free surface. (221) However, experimental evidence indicates that cracks may also form within subgrain boundaries (210, 215,219,220, 222) and that they are first observed using optical microscopy in persistent slip bands, (208,217, 220, 223, 224) that is, in slip bands which fail to disappear on polishing to a depth of several microns. Recent studies have regarded crack growth as occurring in two stages. (207,208, 210, 221) The first stage is assumed to represent a continuation of the shearing process which nucleated the crack. In this stage the cracks follow paths of higher resolved shear stress and are contained within cell boundaries or slip bands. With additional fatigue cycles, cracks which do not terminate tend to propagate in a growth mode which is influenced by normal stresses at the crack tip. (210) In the growth mode preferential propagation through subgrain boundaries occurs. Fatigue crack growth rates have been measured using both optical (210, 244, 252) and x-ray microbeam (245) techniques. The optical measurements on single (210) and polycrystalline (244, 252) aluminium indicate a constant growth rate over a large number of cycles; however, near fracture larger growth rates are observed. (244) The crack growth rate for three single crystals over 50-600 x 103 cycles varied from 4.5 x 10 .8 to 4.0 x 10 .7 cm per cycle (1 800 cpm, constant load), (244) the faster rates being observed in crystals oriented for double slip. But, x-ray microbeam measurements on polycrystalline specimens (245) indicate that the log of the crack length is directly proportional to the number of cycles with higher crack growth rates recorded near crack initiation and fracture. Decreased specimen thickness and testing at reduced pressures (as discussed earlier) result in slower fatigue crack growth rates. (248) Recent 400 kV electron microscopy work (246) on specimens fatigued at room temperature reveals new detail concerning the defect structure in the vicinity of fatigue cracks. It is observed that the substructure is clear and the dislocation density relatively high in the vicinity of cracks. Additionally, voids of sizes 0.1 to 1 x 10.3 mm diameter were observed near cracks, in subgrain boundaries, and in persistent slip bands. Such voids, presumably formed by coalescence from a super-saturated vacancy concentration, probably contribute to crack propagation. Various aspects of fatigue crack propagation have been reviewed by Avery and Backofen, (221) Kennedy, (199) Smith, (222) Low, (205) Thurston, (204) Zackay et al, (206) Meakin and Petch, (228) Beachem, (238) and by Parker and Fegredo. (235) Several other studies of fatigue or cyclic hardening shed insight into the mechanisms of fatigue. The data of Alden and Backofen, (223) obtained by measuring the moment of force needed to deflect constantly aluminium in bending, illustrate fatigue strain hardening as occurring in three stages (Fig. 15a). The initial rapid hardening, followed by a period of constant hardening and finally a period of gradually increasing hardening rate is influenced by strain amplitude and by reducing the temperature to 77 K. These

CRYOGENICS . AUGUST 1972

three stages were correlated with slip line and band development. During the initial rapid hardening stage, randomly spaced slip lines, similar to those found during tensile straining, were observed. Slip band formation corresponded with the third stage, that is, with saturation strain hardening. An interesting speculation, using the electron microscopy findings discussed earlier, is that the rapid, initial hardening may correspond to dislocation multiplication, and subgrain development with subsequent, 'saturation' hardening reflecting contribution by point defects generated during stages two and three. The rate of increase of hardening of all stages is apparently independent of temperature (compare 298 and 77 K, e = 0.002 curves, Fig. 15b), but the extent of stage one hardening is temperature dependent. Finally, the constant strain amplitude tests of Coffin (225) on polycrystalline materials point out a simple, but significant relationship. As plotted in Fig. 15b for commercial purity (99.6 A1) aluminium, the number of cycles (N) to cause fracture are inversely proportional to the square of the constant, imposed plastic strain amplitude (Aep) 1

N~ - -

(6)

(Aep)2

This is significant for two reasons. First, the one-quarter cycle (tensile test) data fall fight in line with the fatigue data. This suggests that fatigue and tensile fracture mechanisms are similar and, as Gilman (226) points out, such fatigue manifestations as intrusions and extrusions should not be expected to play a critical role in fracture. The similarity between low temperature dependence of fatigue strength and tensile flow stress, as discussed earlier, also implies similarities between the rate controlling tensile and fatigue mechanisms. Secondly, since the behaviour shown in Fig. 15b is very closely followed by most materials, it suggests that fatigue life may be increased by decreasing the plastic strain amplitude.

Creep When a constant load or stress is applied to a solid, the solid will.strain in time with the strain rate primarily dependent on the magnitude of the applied stress and on the temperature. The strain of the solid is referred to as creep. If the .applied stress is below approximately the critical resolved shear stress, then only small amounts of anelastic creep (258,259) are observed. Mechanisms of anelastic creep presumably involve localized motion of defects, such as interstitial ordering. On the other hand, if the applied stress is great enough, vacancies migrate and/or dislocations are generated and subsequent creep involves large-scale vacancy or dislocation motion. There are two general types of creep. If the temperature is reasonably low and if the magnitude of the applied stress is near the critical resolved shear stress (or the elastic limit), creep proceeds at a gradually diminishing rate; under these conditions it is eventually possible for creep to cease entirely. Discussions of creep of this type, referred to as logarithmic creep, have been presented. (258260) The more common type of creep proceeds to rupture and has three stages. After the initial strain from loading, the

271

primary stage involves a gradually decreasing creep rate, the second stage corresponds to steady-state conditions in which the creep rate reaches a minimum and remains constant for some time and the third stage features an increasing creep rate, finally culminating in specimen rupture. Most creep experiments and theories pertaining to aluminium have attempted to clarify the second stage in which the rate of defect creation equals the rate of annihilation and 'steady state' conditions are achieved. Attainment of this stage is a function of applied stress,temperature, and material. Creep studies have stimulated a large amount of empirical theories, characterizing the strain, strain rate, stress, and temperature dependence. These formulations have been well summarized by Kennedy. (199) Creep data have recently been reviewed by Bird, Mukherjee, and Dorn (261) while Lubahn and Felgar, (262) Garofalo, (259) Schwope and Jackson, (319) and several seminars (263,264) have also contributed data and review.

Strain-time dependence. A considerable amount of creep data have been accumulated on aluminium, especially by the Sherby-Dorn group. (265-277,320) Typical creep strain-time dependence for various temperatures and stress levels (265,266) are shown in Fig.16. These data when plotted on log scales for compactness and clarity do not dearly delineate the three stages of creep, but are representative of the results of many. (267-289, 325-327) Little systematic creep data at cryogenic temperatures have been reported. Results have been reported for 99.996 A1 at 77 K (274) under a constant stress of 15.4 kg n,m "2 and at 273 K (320) under a stress of 5.6 kg mm "2. Furthermore, activation energies have been obtained (290) for single and polycrystalline specimens at temperatures above 77 K. Single crystal creep tests have been reported by Fleischer and Backofen (291) (0.8-1.0 kg mm'2), at 4.2 K and by Koval et al (292) (99.5 AI, 0.7-3.1 kg mm'2), in the temperature range 1.4 to 4.2 K. Other single crystal creep results, in addition to the above mentioned low temperature data (291,292) have been reported by Schwope, et al (293) (99.99 A1,309 K, 0.140.28 kg mm'2), Johnson, et al (294) (99.99 A1,533-755 K, 0.053-0.21 kg mm'2), Michelitsch (295) (99.99 A1, 295 K), Weertman (296) (99.99 A1,430-889 K, 0.0251.93 kg mm'2), Underwood and Marsh (297) (99.99 A1, 300-866 K, 0.035-0.21 kg mm'2), Voloshina and Rozenburg (324) (293 K, 1.16 kg mm'2; 573 K, 0.124 kg mm'2), and by Rocher et al (328) (76-900 K). Bicrystal results have been reported by Voloshina and Rozenberg (323) (573, 0.20 kg mm'2). The creep strain-time relationship is quite orientation dependent, especially in the temperature range 1.4 to 4.2 K according to the results of Koval et al. (292) The latter work reported normal creep behaviour (greater e for a given time and o) as the test temperature was raised from 1.4 to 4.2 K for [111] and [100] oriented crystals; however their data indicate that, as the temperature is raised (1.4 to 4.2 K), [100] oriented crystals exhibit less strain at constant time and applied stress. Clearly, more low temperature creep studies are needed.

Stress and temperature dependence. The steady state creep rates (2nd stage) which Servi and Grant (278) obtained for commercial purity and high purity aluminium

27_9

id

' 3.SI kg men-2 422K

~-..82kg mm"2 422K

1.97k9 mm-2 4,?.2K

1.41k9 m~2

i~

1.41 kg mm"~ 47"/K ~

1.41kgnm - 4 2 2 K

1~2K::r4 .............................................................. 1 0-3 10-2 iO-I I 0 0 I01 10 2 I0 3

I0 4

Time, h Fig.16

Creep strain o f a l u m i n i u m as a f u n c t i o n o f t i m e . T h e applied stress and test tem~oerature are i n d i c a t e d . Data e x t r a c t e d f r o m S h e r b y a n d D o r n 2 6 5 , 2 6 6 (1.02 kg m m -2 = 0.1 N m "2)

,08 533 K

id

siz~ .~m a'm

366K 1.0-3.0 m m ~ - -

groin

Z

io6

z¢

"~'m~Z"" ,,.. ,,-. " " ~

,..m..ram~m

io°

~ - 756K _-~

no'l~ "

u')

icy ---- 99.3 At 99. 9 9 5 At

Odi(~ I

4

, ,,J*,,,I , ,~,,,.I n ,,*~,.I i ,a~,,.I * ,,,,,. 10-3 I(~ 2 I 0 -I I0 0 IO

Minimum cntep rate, h "° Fig.17 Minimum creep rate as a function of applied stress at indicated temperatures and for indicated purities. Data from Senti and Grant 273

as a function of stress and temperature are presented in Fig.17. From these data, it is quite apparent that an increase in stress and temperature results in higher steadystate creep rates. The dependence of the creep straintime curve on the constant applied stress and temperature are also illustrated in Fig.16. As one would expect, higher applied stresses promote larger strains at constant time and temperature, while higher operating temperatures result in larger strains at constant time and applied stress. Since the creep rate (slope of strain-time curve) normally remains constant over several magnitudes of time, it is customary to consider the effects of the major variables, (stress, temperature) on this constant portion of the creep rate. The stress and temperature dependence are reviewed in more detail by Garofalo. (259) The creep rate is critically dependent on surface condition and on the increasing surface layer stress which impedes dislocation motion. When specimens are electrolytically polished during creep, the creep rate is found to increase by a factor of about 100. (89) It is now commonly accepted that thermally activated processes control the creep rate of metals. The temperature dependence of the process is contained within the exponential term -AH/kT(see equation 7 or 8). If only one process is responsible for the creep process, then the activation enthalpy, A/-/, refers to the thermal energy necessary to activate the controlling process. A number of experiments have been attempted to measure AH, holding such

C R Y O G E N I C S . A U G U S T 1972

variables as o and the internal structure constant during measurement. Usually the change of strain rate has been measured following a sudden change of temperature.

Effects o f purity and grain size. Data (278) presented in Fig. 17 indicate a very significant effect of purity on the creep rate of aIuminium. For example, using the data from Fig.17 for an applied stress of 0.1 kg mm "2 at 644 K the creep rate of 99.995 A1 is found to be approximately eight orders of magnitude greater than the creep rate of 99.3 AI! The same wide divergence of creep rates for the two purities is apparent from the 533 and 756 K data. Grain size influences the creep rate at lower temperatures (for example, at 366 K reduction in grain size from about 2 mm to 0.05 mm average diameter reduces ~ by about one order of magnitude) but at higher temperatures grain size variation has little effect. (280)

Empirical f o r m u l a t i o n

As mentioned earlier, a great many mathematical expressions have been proposed (299) to characterize the dependence of creep strain (e) or strain rate (~) on temperature, time, stress, and some arbitrary parameters. Many of these have resulted from creep experiments using aluminium. The following is a very brief outline of the more promising treatments of general applicability which have evolved during the past few years. The characterization of the minimum, steady-state creep rate (Fig.16) has taken on many slightly different forms during the past twenty years. Earlier formulations of the strain rate were presented in terms of the hyperbolic sine function = K(exp

Ak H T /~ s i n h B °

(7)

where K is a structure-dependent term, B is a structureindependent term, AH is the enthalpy of activation, k is Boltzmann's constant, T is the absolute temperature, and o is the applied stress. The structure dependent K was found to decrease with increasing pre-strain (dislocation density). (272) The data also reveal that lIB is linearly proportional to the atomic percent solute alloying element, but is insensitive to changes in defect structure. (272) Although as of this date there appears to be no unanimity, several expressions conform to much of the available data. Weertman (258) and Bird et al (261) for aluminium suggest that the stress-dependent strain rate (e0) is proportional to o n at low stresses where n -~ 4.5 and o is the applied stress. The general form of this expression is ~

=

(exp

AH~co2

- kT ]

[A°~2"5

sinhk~-)

(8)

where C and A are constants, the latter corresponding in concept to activation area. At low stress (8) reduces to a simple power law relationship. A slightly different function has been suggested (298) by Garofalo for the stress dependence of the minimum creep rate = C sin h (ao) n

CRYOGENICS . AUGUST 1972

with C and a being constants at a given temperature. This same formulation has been proposed to characterize the hot working behaviour of metals. (291) This expression, which also reduces to a power law expression (e0 = c'°n) at low stress levels, approximates the exponential expression, e0 = C" e B°, at higher stress levels and near the melting point (at low stresses) satisfies the linear expression, eO = C'"o. Garofalo (259) computes values for the materiai parameters C, ct, and n as follows: at 477 K, C = 2.8 x 10 -6 s "l , tx = 5.6 x 10 -3 cm 2 kg "1 , n = 5.0, at 533 K, C = 1.9 x 10 -5 s"l , ct = 7.3 x 10-3 cm 2 kg"1 , n = 4.55 and at 920 K, C ---2.7 x 10 .8 sl , t~ = 1.8 x 10 -3 cm 2 kg l , n = 1.24. Defect structure

Fewer electron microscopy studies have been performed to determine aluminium defect structures in creep than in the case of fatigue. However, using results mainly on other materials, Garofalo (259) and Bird, et al (261) have provided excellent reviews of research progress on this subject. It appears that dislocations generated during the primary creep stage gradually form a subgrain structure. During steady state creep, this subgrain structure remains relatively stable (within the limits of present day measurement techniques). Substructure evolves faster and the average subgrain diameter is smaller at higher temperatures and stresses. Typical subgrain diameters obtained from optical microscopy measurements for aluminium during creep have been reported (300) as 0.030 to 0.5 mm within the temperature range 533 to about 900 K; however, it has also been argued that these are too large. (261) Furthermore, these sizes are considerably larger than those formed during tensile or fatigue testing (see Fig.21). Electron microscopy data (321) indicate a decrease from 0.060 to about 0.004 mm subgrain size on creep of 99.9995 A1 under a constant stress of 0.7 kg mm "2 at 523 K; these values seem more consistent with other property investigations. While a low dislocation density is present within the subgrains, a large concentration of dislocation loops has been observed near subgrain walls. (309) There is evidence that these subgrain boundaries have considerable influence on the creep rate, the smaller subgrain sizes and more stable boundaries promote reduced rates. (309) Earlier studies (276, 277,294, 300-307) established that both coarse and fine slip occurred during creep and optical microscopy and x-ray techniques revealed that a substructure existed. Apparently, the average subgrain diameter is less than the average slip line spacing, (259,308) although surface slip line resolution using optical microscopy conceivably could be responsible for the discrepancy between the two sizes. Grain boundary sliding has been found to contribute approximately 1 to 20% of the creep strain. (277,295, 310, 311) Maximum contributions from sliding seem to occur at low creep strains and high temperatures.

Defect mechanisms

(9)

Primarily by using the inferential activation energy, area, and volume data, it is possible to speculate on defect mechanisms controlling creep rates. As many as five distinct defect mechanisms may control the steady state creep of aluminium at various temperatures: (1) dislocation glide,

273

(2) dislocation cross slip, (3) dish)cation climb (subgrain formation), (4) grain boundary migration, and (5) mass vacancy diffusion. These mechanisms have been adequately discussed by Schock. (317) Mechanisms (4) and (5) contribute significant creep strains only at temperatures close to the melting point; mechanism (3), being dependent on self-diffusion, is significant at temperatures higher than about 400 K (climb appears to be the predominant mechanism above 600 K), while dislocation glide and/or cross slip contributes to creep strains at temperatures below 600 K. A reminder is needed that the above temperature ranges pertain to steady state creep (constant ~ for a period of time), since at all temperatures primary creep is thought to be controlled entirely by dislocation glide. Since the experimentally-determined activation energy (1.54 eV) is constant from about 600 to 900 K for polycrystalline specimens and since this value corresponds to the energy for self-diffusion, the conclusion is often reached that the dominant creep process in this temperature range is dislocation climb. Apparent activation volume values are also constant in this temperature range and indicate that the process proceeds over slightly larger than one atomic volume. In the temperature interval 260 to 400 K the activation-energy has a constant value of about 1.21 eV for polycrystalline specimens, while the activation volume (slightly less than 1 atomic volume) remains constant from about 350 to 420 K. The activation energy corresponds closely to that estimated for cross-slip (recombination of a screw dislocation over 10 atomic distances) of 1.13 eV. The low temperature process shown by Kramer (84) to have a low activation energy in the vicinity of 0.03 eV, but very dependent on surface conditions, could very well correspond to a Peierls-type energy. (292) Dislocation intersection interactions are thought to require higher energies of the order of 0.48 eV. (317) It has recently been argued in a series of papers, (312-315) that the steady state creep rate is controlled by the balancing of internal stresses (oi) and recovery. Internal stresses are built up by the accumulation of dislocations during glide in the primary stage. Recovery acts to relieve such stresses, presumably by dislocation climb. The effective stress (°e), then, to maintain the creep rate may be simply expressed as

oe = o - oi

(1 O)

with o the applied stress. In the absence of recovery o i would increase to o, and creep would cease. This approach,

-E E

6O Guillct. Cournot 0922)-99.7At 50

.o 40 E'~

o .z3 E E

0

30

o~

26

--

8u~

24

~'~

22

E

20 18

.O E

16

IE

,4

NBS data unpublished,

"r

BGttler (19531 ', Helling,Ncunzig 119511 Hikog¢ (,053)

T

1[/

i

,#~Ox /

~

Guill~, CournotllO2a} . ('O00kg, lOmm ball) i Tomlin$on(BS7)

T =

Ti ~ / "

II, I

~ ~'/

[/X

~. -"

~

El

E

r-

=

~

L

8,o-'

rn

i

I

I

I

I

ioo

lo,

jo2

loa

lo4

iOs

Aluminium impurity, ppm

F,ig.19 Room temperature hardness of aluminium as a function of random impurities

although not new in concept, should stimulate new research pertaining to internal stress measurement, which has already been reported in aluminium. (312,316) These measurements indicate that o i / o decreases with increasing o and T with an abrupt change in temperature dependence occurring at about 460 K. At present there is no unanimity regarding this approach and theories relating creep rate solely to recovery or to jogged dislocation motion still have a respectable amount of support. Creep rupture in aluminium is thought to result from grain boundary void (or cavity) growth. Comparatively little research has been reported on creep fracture in aluminium; most evidence, obtained from optical microscopy, indicates that voids play an important role in fracture propagation.

Hardness Most commercial aluminium hardness data are presented as Brinell hardness numbers in kg mm "2 units, obtained using 500 kg loads applied to a l0 mm diameter ball. These numbers range from 18 to 23 kg mm "2 for a purity of about 99.99 A1 at room temperature. Typical temperature-dependent data are shown in Fig.18. The values for most of these investigations (330-332) were converted assuming a linear relationship between quoted hardness numbers and Brinell hardness numbers (500 kg/ 10 mm). Despite the time span of over fifty years, there is a good agreement between the data of Petty (330) and Kurth. (332) Recently, the relative variation of hardness for annealed and cold worked aluminium subjected to a radiation was measured (333); increases up to 1.5 mm "2 were obtained with maximum radiation.

We (334) have determined the room temperature DPH hardness under 200 g loads for a large range of purities. Purity ranged from 99.0 A1 to some of the highest purity ,° aluminium ever. manufactured. The data, shown in Fig.19 ......... ~ ~ Ivonko(19681 together with that reported by others (329,335-341) Kimulra' Nalkna ° ( 11971)" i i ~ 1 I 0 indicate that the hardness may be insensitive to purity 0 I00 200 300 400 500 600 700 800 900 I000 when the impurity level is less than about 100 ppm. Our Temperature, K measurements, obtained within individual grains, do not reflect possible grain boundary contributions, Impurity Temperature dependence of hardness of aluminium Fig.18 ~

274

20

C R Y O G E N I C S . A U G U S T 1972

levels were obtained using 4 K electrical resistivity measurements and assuming that on the average 1 ppm impurity residual in zone refined aluminium contributes a resistivity of 0.6 x 10 -9 ~ cm. (1) The effects of small binary solute additions of V, Nb, Ta, Cr, and Mo on the hardness of aluminium have recently been reported (338) as a function of temperature (293-673 K). Additions of concentrations as low as 0.1% of these elements contributed to a noticeable hardness increase.

for which tensile strengths in excess of 170 000 psi have been reported! Yet, there are other ways to increase the strength: for example, the introduction of excessive dislocations, through subgrain formation or by rapid quenching techniques, serves to harden aluminium. Smaller grain sizes also result in higher flow stresses. (The effect of grain size was previously discussed.) Electron or neutron irradiation, through the introduction of point defects, significantly increases the yield strength in the temperature region below which these defects are mobile. Additionally, below about 150 K, thermal fluctuations are decreased sufficiently to permit very local stress fields to become obstacles to dislocation motion.

A recent compilation of hardness data has been presented by lvarlKo. (339) This handbook contains Russian aluminium hardness data published from 1921 to 1960 which range from 15 to 25 kg mm "2 for annealed specimens at room temperature. Additional data on hardness versus applied load are presented by Upit and Varchenya. (340) Their results (339,340) are conflicting, but normally it is thought that larger loads result in larger hardness values.

All methods of strengthening serve to decrease dislocation mobility, acting as obstacles to their motion under applied stress. Some obstacles, such as grain boundaries and precipitates have large stress fields. Their contribution to strengthening is thought to be more or less independent of temperature (referred to as athermal mechanisms). Other obstacles, such as point defects, have smaller associated stress fields; thus dislocation motion past these obstacles may be substantially aided by thermal energy and their strengthening contribution is only observed at low temperatures. An excellent review of thermally-activated strengthening mechanisms has been recently presented by Nix and Menezes. (346)

Impact Apparently only low temperature impact tests on highpurity or commercial purity aluminium have been performed. (144, 343-345) These results are presented in Fig.20. Absence of higher temperature data is not surprising, however, as the ductility of aluminium causes specimens to bend rather than fracture at room temperature. (144) At higher temperatures aluminium impact specimens would most certainly not fracture, making fracture energyadsorbed data unobtainable with conventional testing techniques and specimen geometries. Low temperature data indicate gradually increasing energy to fracture down to 76 K, followed by a slight reduction in strength at 20 K. It is safe to conch,de that aluminium, under impact loads, is tougher at cryogenic temperatures than at room temperature.

The major techniques of strenthening are reviewed below. These discussions are necessarily cursory, with only limited discussion, explanation, and references being included.

Alloying

When foreign atoms are added to aluminium, two types of strenthening may occur: solid or precipitation. Each has been the object of extensive research and good reviews exist of solid solution strengthening (347-351) and of precipitation hardening. (351-354)

Strengthening

Data on solute strengthening effects exist for most aluminium binary alloy systems. The most recent work by Tensi et al (355) on A1-Mg single crystals also contains an adequate reference list of past studies.

Aluminium, although relatively soft in the annealed condition, is a comparatively easy metal to strengthen. In the late 1920s tensile strengths as high as 50 000 psi rib in "2) were achieved in aluminium by alloying and cold working. Today, using combinations of cold working, alloying, and precipitation hardening, tensile strengths up to 100 000 psi have been achieved. Yet more impressive, in the past five years an almost quantum jump in strength has been obtained in material composed of boron filaments in aluminium alloys

80 -

Fontana(19481 99.6 At, Charpy V

..m

60 W a r r e n ,

Reed (1963)- 99.6 A I, I/2 Charpy V

40

Colbeck,Mac Gillivmy (i933)99.6 At, Izod

IOO

"3 ~n

so

~.o 20 i

0

Johnston, Brooks (1952)-99.6At, Chorpy Keyhole t i t , i i • i | t i i i i

_

IOO 200 300 400 5CXD 600 700 800 Temperature,K

Fig.20 Variation of impact energy adsorbed as a function of temperature - type of impact specimen is indicated

CRYOGENICS. AUGUST 1972

o

The following single crystal trends have been documented by Tensi et al (355) for A1 -Mg "alloys.As solute concentration is increased, r 0 and 1"111linearly increase, 7"1i/70 increases very gradually, 01 increases abruptly to concentrations of about 0.I at % Mg, then slowly decreases, 01i decreases linearly in the concentration range of 0.1 to 2.0 at %Mg and the strain at the beginning of stage II increases approximately linearly with alloy addition. The linear dependence of r 0 on Mg concentration is also found in the A1-Ag, A1-Cu, and perhaps in the A1-Zn system; data which indicate steeper 1"0 dependence on concentration have been reported for the A1-Zn system. In Table 3 plots are presented of the slopes of 1"0 versus concentration for A1-Zn, Cu, Mg, and Ce alloys assuming a linear dependence. Table 3 also contains incremental resistivity data that may be expected per 1%. These are included to provide relative information with respect to resistivity increases which may be expected for various hardening mechanisms information which is vital for appropriate design of cryogenic conductors. These conductors, in general, need to have good strength but low resistivity. Notice that, for the alloying element included, Ce apparently provides

275

the best strength increase at the expense of the accompanying resistivity increase. The flow stress of polycrystalline aluminium alloys is almost linearly dependent on solute concentration. (356, 357) At lower temperatures the linearity is well marked, but at temperatures above approximately 200 K the dependence is slightly non-linear with higher concentration additions contributing slightly less to the increase of flow stress. Table 3 contains approximate values for the increase in flow stress per 1% alloying additions in solid solution of Zn, Mg, and Ce. The Ce data, relying solely on one source, (358, 359) indicates an abnormally strong increase for this element, compared to the other elements. From the Zn, Cu, and Mg information it is apparent that considerable differences exist in the solid solution strengthening effectiveness of binary element additions. Solid solution hardening is attributed to three effects. (351) (1) The addition of foreign solute atoms changes the

electronic and chemical nature of the aluminium matrix which, in turn, lowers the matrix stacking fault energy. A lower stacking fault energy results in a wider than average separation of partial dislocations which make their recombination more difficult during cross-slip or jog formation processes. (2) The foreign atoms, usually of differing atomic volume from that of aluminium, serve to internally strain the matrix. This internal strain provides increased frictional stress on moving dislocations, serving to retard their motion. (3) Dislocations act as partial sinks for foreign atoms, providing lattice sites in which the foreign atom stress fields arc reduced. However, atmospheres of foreign atoms about dislocations act as locks, making dislocation motion considerably more difficult. All three of these effects serve to increase the hardness and strength of the alloy compared to pure aluminium. Precipitates usually increase the strength of the aluminium alloy considerably more than does solid solution-strengthening. For example, the critical resolved shear stress of A1-1-7

Table 3. Strengthening and resistivity contributions of various types of defects

Type of defect Dislocation loops

Dislocations

Technique

Single crystal, quench from 870 to 263 K; 300 A diam, 10.0 x 1014 cm-3 Cold work using assume 1010 dislocations

r = o~GbN ~ ,

Strengthening increase (g mm "2 per defect) 9 x 10 "14

Approx recovery temp

Resistivity increase ( ~ cm per defect)

~45 K

6.6 x 10 .24 cm 3

1 x 1010

~500 K

3 x 10 "19 cm 2

1 x 1012

~500 K

16 x 10 .9

2 x 1012

Ap \ [2 cm ]

Ar0 loop cm "3 4 x 10 .7

Ar line cm "2

produced, a = 0.5 Cold work (76 K) Single crystal Vacancy, interstitial

3 800 e = 0.6*

Ar - -

Electron irradiation 550 (~6 x 1017 elect cm "2 A r = 6 x 1017 elect cm "2 A r = A C '/= , at 20 K)

e = 0.6 ~<250 K

7 x 10 "10

8 x 1011

6 x 1017 elect cm "2

A//owng

Zn

167

Solid soln (<3 at %)

Melting pt 0.23 x 10 .8

ATO - _ _

1 at%

1 at% Hardened ( 2 - 3 at %)

3 900

7 x 10 8

Melting pt 0.23 x 10 -6

2 x 1010

Ar0 -

1 at%

1 at% Cu

1 300

Solid soln (~<3 at %) Ar o -

Mg

1 at% 600

Solid soln (~<2.0 at %)

Melting pt 0.83 x 10 .6

2 x 109

1 at% Melting pt 0.46 x 10 -6

1 x 109

Ar c -

1 at% Ce

3 000

Solid soln

1 at% Melting pt 0.066 x 10 .6

5 x 1010

Ar1 at%

1 at%

*Dislocation density not known, only resistivity and stress increases per unit strain.

276

CRYOGENICS . AUGUST 1972

at % Cu in solid solution is 3.5 kg mm "2 at 273 K but may be as high as 10.5 kg mm -2 if the alloy is optimally aged or precipitation hardened. (360) The AI-Zn alloys demonstrate similar properties: an AI-3.5 atom % Zn solid solution alloy has a r 0 value of 0.7 kg mm "2 which increased to 4.5 kg mm "2 after precipitation (361) (see Table 3 for other A1-Zn data). Similar strength increases are observed in preciptation-hardened polycrystalline aluminium alloys. Basically, preciptation creates two types of barriers for dislocations. (1) Precipitates (or, alternatively, ordering or clustering of atoms in pre-precipitation zones) of a second phase (differing crystal structure) provide lattice discontinuities and stress fields which dislocations must circumvent during slip. Number and size of particles, particle crystal structure and morphology, and particle-matrix coherency all contribute to the magnitude of the stress field and thus to dislocation immobility. It has been shown (362) that in A1-Cu single crystals r 0 is approximately linearly proportional to the reciprocal of the average precipitate spacing. (2) Precipitation tends to begin in the vicinity of dislocations. Presence of second phase particles within dislocation stress fields or cores serves to reduce the number of mobile dislocations.

Work

hardening

From the time that man first began to forge weapons and tools, it has been known that permanent deformation of metals results in increased strength and hardness. It is therefore not surprising that the earliest commercial alloys were partially cold worked, resulting in increased strength. Many papers have been concerned with work hardening theory and experiment. Particularly, single crystal studies (363,364) and 01 and OII interpretation have received much emphasis. For discussion and other references see Conrad, (6) Nix and Menezes, (346) Parker, (365) and the very good work hardening conference proceedings. (366, 367,368) Apparently, only scattered empirical studies (usually performed by aluminium company research groups), have been directed toward maximizing the combined effects of deformation temperature, recovery, and alloy composition on strengthening. Instead, a large amount of work has been devoted to room temperature work hardening of specific commercial alloys. One isolated study implies, however, that greater retained strength can be achieved by deformation at temperatures lower than room temperature. (370) Strain hardening results from the generation of lattice defects, predominantly dislocations, during plastic deformation. Interactions between dislocations serve to decrease their mobility, so that higher applied stresses are requried to move them and to cause plastic deformation. Specifically, it is considered that the applied shear stress, r, is proportional to the square root of the dislocation density, N (equation 3) and the resultant strain rate, ~, is related to the mobile dislocation density, Nm, and the average dislocation velocity, v by the expression

=NmbP

(ll)

The relation r = o~GbN V2(equation 3) has.been approximately experimentally verified using x-ray topography,

CRYOGENICS

. AUGUST

1972

(370) electron microscopy, (372) and x-ray line broadening. (373) The measured et values varied considerably: tension of single crystals at temperatures between 433 and 693 K monitored by transmission x-ray topograph produced an a value of about 0.77, (371) tension of polycrystalline specimens at room temperature monitored by electron microscopy produced an ot value of 0.35, (372) high temperature (623 K) tension of polycrystalline AI, using electron microscopy to judge the dislocation density, resulted in an a value of 0.018, (373) and x-ray line broadening analysis of crystals strained in tension at room temperature produced an t~ value of 0.06. (273) In Table 3 an ~ value of 0.5 was assumed, together with a dislocation production of 1 x 1010 dislocations per cm 2 and the validity of (3). The resultant stress increase per resistivity increase is better than that obtained from solute hardening. Better still is the single crystal deformation at 76 K hardening-resistivity ratio (also in Table 3); indeed, this combination produced the best strengthening per resistivity increase of any combination analysed and presented in Table 3. Su bstructu

re

Plastic deformation of aluminium is controlled by the generation and motion of dislocations. Most, if not all, dislocations are initially both created and move on [ 111 ] planes at room temperature. However, cross-slip and climb are easy in aluminium due to a low stacking fault energy so that even at very low strains dislocations tend to move out of their glide planes and cluster into tangles, creating regions of relatively high dislocation density. During subsequent deformation or during recovery, these regions reform into a cell or subgrain structure. The average cell size decreases with increased plastic deformation as more and more dislocations become trapped in the cellular structure. Fig.21 presents available x-ray and electron Fatigue cycles

Joa 64X

1(53

,o"

i i

,

I ,

IO

20

30

Ic:; ,

io 6

i ,

,

50

60

, i

io 7 ,

i i

56 48 40 E E

32 24

U

16

O

40

70

80

90

IOO

Deformation, Oto Fig.21 Dislocation substructure size resulting from plastic deformation. Curve A -- Kelly (1954), tensile at 298 K; Curve B -- Swann (1963), rolled at 298 K; Curve C -- Swann (1963), tensile at 77 K; Curve D - Grosskreutz (1963), fatigue at 298 K, use upper scale (fatigue circles); Curve E -- Stein (1967), creep at 533 K, 1 000 psi; Curve F -- McLean (1952--53), creep at 473 K

277

microscopy measurements (301,212, 321,374, 375) on the cell size as related to degree of plastic deformation (or number of fatigue cycles) for the temperatures indicated in the figure. The measurements are reasonably consistent (except for the fatigue results) and indicate that a cell size of 1 to 3 x 10"~ mm is rather quickly formed within strains of 0.10, and that this size remains constant throughout subsequent deformation. Exact correlation of this substructure with optical microscopy observations of slip lines is not clear at this time. Also unresolved is the responsible mechanism(s) of dislocation glibe and/or cell wall motion to produce plastic deformation.

who obtained flow stress dependence on both grain size and substructure for the same specimens. At present it is not clear why substructure is a more effective strengthening agent. It may be that grain boundaries act primarily as an emitter of dislocations, contributing to hardening by increasing the dislocation density; whereas subboundaries act as barriers to dislocation motion, contributing to strengthening by creating long-range stress barriers.

Irradiation Throughout this review several brief references have been made to irradiation effects on the deformation properties of aluminium. Irradiation strengthening can only occur at temperatures below which point defects or point defect clusters remain relatively immobile, that is below room temperature for aluminium. Therefore, from an engineering standpoint it is not practical to consider electron, neutron, or alpha particle irradiation as a means of strengthening aluminium if the final product is warmed to room temperature. However, the effects of irradiation do serve to illustrate the considerable strength increases one can achieve by point defect introduction in the lattice and allow studies to be conducted under controlled conditions. General reviews of irradiation effects have been published (379-382) and at least two major conferences (383,384) have been held on this subject lately.

Flow stress is critically dependent on substructure size. Fig.22 shows the room temperature dependence of tensile flow stress (57, 58, 94, 109, 112, 376) and of hardness (342, 377) on subgrain diameter. These data, while not conclusive, indicate that the critical resolved shear stress of single crystals is less dependent on subgrain size than is polycrystalline flow stress. A distinction must be made between the substructure created by plastic deformation (Fig.21) and the substructures considered in the measurements reported in Fig.22. The latter substructures were created by small deformations (e -~ 0.05 - 0.20) at room temperature followed by holding at higher temperatures. This procedure produces cleaner substructure walls (less tangles, more polygonization) and probably results in segregation of impurities to the walls. (378) Such segregation should produce a pinning effect which may help explain the exceptionally strong flow stress dependence on such substructures.

Several investigations have studied electron (47, 73, 82, 385,386) or neutron (86) irradiation effects on aluminium single crystals. The primary effect of irradiation is to increase r 0 or the yield stress with this increase in flow stress being maintained during subsequent straining. A large increase in r 0 from about 0.2 to 3.0 k~ mm "2 for an integrated electron flux of 2.8 x 1017 e cm "L has been reported. (385) Flow stress increases have been shown (386) to vary

Comparison of grain boundary strengthening (Fig. 10) with subgrain strengthening (Fig.22) indicates that the latter has by far the larger effect on the flow stress. This was nicely shown in the experiments by Hultgren, (114)

as

16

Ar ~ ~/2

(12)

14

12 D

t~E

IO

E

t--

6

"l-

4

2 os

AG,, id 4

j6 3

.,,<.~'B 16a

161

io °

IO I O

Subgmin diameter, mm Fig.22 Flow stress or hardness resulting f r o m variation o f subgrain diameter. Presumably grain size is constant. Curve A Hultgren (1964), gs = 0.3 mm, e = 0.0025 in tension; Curve A' -gs = 0.13 ram; Curve A " -- gs = 0.04 mm; Curve A ' " -- gs = 0.016 mm; Curve B -- Lake, Craig (1968), 70; Curve B' - Lake, Craig (1971) r0; Curve C -- Ball (1957), proportional limit; Curve D - Hockett, McQueen (1970), c = 0.002 in tension; Curve E - Wong et al (1967), hardness; Curve F -- Cotner, Tegart (1969), hardness; Curve G - Hansen (1969), e = 0.002 in tension, 0.2% A I 2 0 3

278

where • is the integrated electron flux. The number of defects produced is proportional to this flux. The strain during easy glide in stage I may be slightly increased by irradiation. The strain rate sensitivity of the flow stress (At/AinU,) was found (385) to increase proportionally with the square root of the electron dosage. Polycrystalline stress-strain measurements after neutron irradiation have also been reported. (181,182) Significant aspects of their results have already been discussed. Large decreases in elongation were observed after irradiation. Table 3 presents electron irradiation stress and resistivity increases for comparative purposes. Two major factors make irradiation hardening impractical to consider seriously: (1) reactor design and particle penetration depths place severe limitations on material size and cross section and (2) considerable recovery, that is loss of strength, begins at about 60 K upon warming. Recovery after irradiation is discussed more comprehensively in the section on recovery.

Quenching

In 1955 Madden and Cottrell (63) noticed that when aluminium is quenched, the flow stress is increased over that

C R Y O G EN ICS . A U G U S T

1972

which is slowly cooled. Since then, the types of defects introduced by quenching and the resultant effect on flow stress has been studied. This subject, being relatively new, has not been adequately reviewed. As quench hardening has possible practical applications, more details will be presented here. Defects observable in the electron microscope form from condensation of excess vacancies during quenching of aluminium from temperatures above about 700 K. There is evidence, also, that vacancy clusters not observable using electron microscopy may be formed by quenching from temperatures as low as 630 K. (145) The vacancy fraction Nv/N, where N is the number of available lattice sites and N v is the number of vacant sites, is Nv

-E v

- exp--

N

(13)

kT

where k is Boltzmann's constant, T is the absolute temperature, and E v is the vacancy formation energy. For aluminium E v is about 0.76 eV. Typical quenching temperatures range from 750 to 900 K, with quenching bath temperatures ranging from 200 to 298 K. During quenching and during ageing within the temperature range 200 to 300 K, vacancies move to reduce their concentration through annihilation or coalescence. Vacancy motion is proportional to

3000

Edl;'~i:t;(i i '0%ilt)ot~9~~6,ii i: ¢;~i:: q Cottcrill (1063),to 273K,sesslle~/

2 000

o<

Yoshldo(19¢x3),from Yoshido(1903),to ~ ~/ 823K, s.ssil. 2"L3K,s.ssil. ~ ' N N ~ "~Cotterill(1963),to 273K,sessil¢ "~ •

N

1000

2O0

a

200

iO Is

u

800

1000

Yoshida{1963)to 273Kprismotic

~

ottcrill(1963)to 273Kprismotic~

iO14

Edington (1965)to 296K$¢s$11¢ Cottcrill (19631to 273Ksessile

iO13 Yoshido(1963), from 823K,sessile iO12

Cott,rill (1063).to 273K.pp;:n~tt:: Edington (Ig6S),to 296K 200

b

600 Temperature, K

'Eu e-

400

6o--o

400

8

oo--

IOOO

Tcrnpcmture K

Fig.23 Effectsof pre-quenchannealingtemperatureon quenchedin dislocationloop size (a) and loop concentration (b) CRYOGENICS . A U G U S T 1972

exp - (E v + Em)/!cT, where E m is the energy of motion and is equal to about 0.7 eV. Typical site jump frequencies for aluminium vacancies are about 104 s"1 at 873 K. (387) The type of defect formed seems to be critically dependent on the holding temperature, on the rate of quenching, on the temperature of the quench bath, and on the subsequent ageing time and temperature. A multitude of quenched-in defects have been identified: prismatic dislocation loops (387-395,414) faulted (Frank) sessile dislocation loops, (392-403,414) two layer sessile dislocation loops, (393, 399,403,414) voids, (144, 404-408,414) two triangle two-layer loops, (399) three-layer dislocation loops, (400) four-layer loops, (402) dislocation climb sources, (399) and local clusters of unknown geometry. (145) Loops begin to anneal out upon warming to the temperature range 380 to 460 K with annihilitation being complete at about 475 K. Temperature dependence of percent recovery of loop-induced strengthening has been reported. (391,394, 398, 403,413,415,416) Higher material purity reduces the concentration and increases the average size of prismatic dislocation loops for identical quenching and ageing schedules. (393,394, 403) However, evidence of the effect of purity on the concentration of faulted, sessile dislocation loops is conflicting. (393, 394, 403) The results of Edington and Smallman (393) and Cotterill and Segall (394) strongly suggest that it is harder to obtain faulted loops in lower purity specimens. Yet the data of Yoshida, et al, (403) covering four different purities, unmistakably indicate a trend to higher fault loop concentrations in the lower purity specimens. Repeated quenches have been shown to increase the fraction of faulted loops. (393,394) Fig.23 relates the average loop density and size to the quenching temperature and bath temperature. (144,393,394,403) Notice the general agreement that higher quenching temperatures promote larger loop concentration and smaller loop size. Just recently it has been shown (420) that prequench annealing time influences the post-quench number of loops; longer annealing times result in fewer loops. Fig.24 compares the ratio of the flow stress of quenched specimens to that of annealed specimens (each tested at the same temperature) as a function of temperature. (48, 144,410) It is striking that three distinct groups are apparent: one group encompasses flow stresses increased by a factor of 2 to 4, a second group indicates increases by a factor of 6 to 10, while the third group encompasses the largest increases of a factor of 14 to 20. Careful perusal will show that these groups cannot be correlated with loop density nor loop size. Additionally, Westmacott (147) has obtained temperature dependent flow stress data as a function of pre-quench annealing temperatures as low as 627 K to 873 K. Strengthening increases were achieved from all annealing temperatures, even though quenchedin defects could not be observed using electron microscopy in specimens quenched from temperatures below 723 K. It has been suggested (147) that such observations strongly imply that vacancy cluster complexes of very small sizes contribute significantly to strengthening. Also, a severe dependence on ageing temperature is revealed by measurements (415, 416) of flow stress after quenching from a series of temperature (748 to 473 K) and subsequently ageing at a series of temperatures (200 to 473 K). This profound dependence of ageing temperature implies that cluster size is critical. A small ageing temperature range,

279

identified about 238 K for two different purities and independent of pre-quench temperature, (416) produce the largest increases of flow stress. Also, at constant ageing temperature, the flow stress was higher for higher prequench temperatures. Apparently, a given ageing temperature defines an equilibrium cluster size while higher prequench temperatures create more clusters but do not significantly change the average cluster size. Mori and Meshi (418) have shown that deformation occurred in distinct slip bands which they observed were channels, free of vacancy loops. This channelling effect, together with the observed lower rates of work hardening in quenched specimens, provide additional arguments in favour of quenched-in obstacles to dislocation motion. The other intriguing aspect of data presented in Fig.24 is the variation of the temperature dependence of the hardening mechanism(s). While both of the highest strength quenched materials (curves 2, 3) exhibit similar temperature dependence, the lower strength materials (curves 4-9) are, 20

18 (3) (2)

14

12

I0

8

to a first approximation, independent of temperature and the middle strength range (curves 1, 10-12) have mixed temperature dependences. The tendency of the high and low ranges to have internally consistent, but differing temperature dependences, fortifies the contention that different hardening mechanisms are responsible for the three strength ranges. Tanner (411) measured the temperature dependence of the flow stress of quenched aluminium crystals by the Cottrell-Stokes method (91) and found identical behaviour to that of annealed crystals; that is, the ratio of the flow stresses for any two temperatures was constant, independent of strain, and equal to the annealed crystal ratio. Birmbaum (412) subsequently has interpreted these results as confirmation that the short range stress field interactions (identified as r u in equation 1) are associated with the same obstacles. However, results presented in Fig.24 strongly imply that the ratio would not be constant if the temperature cycling experiments had been perf6rmed on crystals strengthened by over a factor of 6. Therefore, it seems that for better understanding of the temperature dependence of the flow stress (equation 1) quench hardened crystals having hardening ratios of at least six should be measured. The temperature dependence presented in Fig.24 is unreliable since no attempt was made to hold the strain-induced defect structure constant. Another interesting aspect of loop hardening is contained in Table 3. The strengthening increment per loop introduced is about 7 orders of magnitude less than the strengthening increment obtained by the introduction of a dislocation through plastic deformation. These data, of course, assume a direct relation between number of observable loops and hardening, which, as we have pointed out, is not always the case.

02) Composites flo)

6

4

2 (8) (4)

O

I

IOO

I

200

I

3OO

Ternpcmtur¢, K Fig.24 Ratio of f l o w stress after quenching to f l o w stress after slow cooling, as a function of temperature. Curve (1) -- Deguchi, Yoshida (1966), 9 9 . 9 9 AI crystal near [ 1 2 3 ] , quenched 913 K to 268 K, loop density = 0.6 x 10 TM cm -3, loop size = 1 0 0 0 ~t, arc at 300 K = 0.055 kg mm-3; Curve (2) -- same as (1) except quenched 913 K to 273 K, loop density = 2 x 1014 cm-3 loop size = 700 A ; Curve 3 -- Same as (1) except quenched 873 K to 263 K, loop density = 10 x 10 TM cm "3, loop size = 3 0 0 ~ ; Curve (4) -Westmacott (1966), 9 9 . 9 9 AI polycrystals, quenched 823 K to 273 K, aged at least 24 hours at 273 K, loop size = 350 A , a a c at 300 K = 0.9 kg mm "2, a at e = 0.001 ; Curve (5) - Same as (4) except quenched 773 K to 273 K, loop size = 1 000 A ; Curve (6) -Same as (4) except quenched 880 K to 273 K, loop size = 2 3 0 / ~ , loop density = 4 x 10 TM cm'3; Curve (7) -- Same as (4) except quenched 928 K to 273 K, loop size = 30 fl~; Curve (8) - Shiotani, Kimura, Hasiguti, Maddin (1967), 9 9 . 9 9 9 AI quenched 673 K to 271 K, aac at 300 K = 0.25 kg mm "2, proportional limit; Curve (9) -Same as (8) except aged at 298 K - 36 h; Curve (10) - Same as (8) except quenched 813 K to 271 K, aged at 298 K -- 36 h; Curve (11 ) -Same as (8) except quenched 873 K to 271 K; Curve (12) - Same as (8) except quenched 873 K to 271 K, aged at 298 K -- 36 h

280

Through combination of two or more dissimilar materials, properties which differ drastically from that of the base materials can be achieved. For example, aluminium alloy tensile strengths can be doubled by the proper imbedding of boron or graphite fibres into the aluminium matrix. During the past ten years, tremendous research and development of such composites has taken place with the primary objective of designing materials for specific job requirements. As a result composites have been produced having higher strength-to-weight, higher modulus-to-weight, and improved fatigue and creep properties compared to conventional materials for many applications. The present discussion will briefly review composite properties and research which relate to an aluminium matrix base. Discussions containing much greater detail have been presented on composites (419-423) and on dispersion hardening, (424, 425) and from conferences. (426) (These above reference lists are not comprehensive.) Two types of metal base composites are currently being examined: those using fibres for strengthening and those which employ particles. Common aluminium base fibre composites are Al-stainless steel, Al-boron, M-beryllium, and M-silicon carbide. Particles which have been combined with aluminium include A120 3 and silicon carbide. Composites of hard particles in an aluminium matrix have been rather well characterized. Particularly aluminium-

CRYOGENICS

. AUGUST

1972

alumina (A1203) particle composites have been examined rather completely. Tensile (109,427-431) and creep (432) properties have been systematically examined as a function of A1203 concentration. Limited impact analysis (433) has been performed. Tensile data (430) for the temperature range 1.3 to 300 K have been obtained for A1 14% A1203, between 76 and 473 K for A1-4.5 to 11.7% A1203, (431) and at 298 and 673 K (109,427,428) for a range of AI203 concentrations. Hansen has quantified his data by the expression

o = o0 + kd -V2

(14)

relating the flow stress (o) to a constant o 0, which is linearly proportional to the A1203 volume fraction, and to the grain size (or subgrain) average diameter d. The constant k is independent of oxide concentration. Large flow stress increases can be obtained using A1203 particles. For instance the addition of 3.8 vol % A1203 at 10% strain increased the flow stress from 7.5 to 21 kg mm "1 at room temperature. (427) At 4 K the 0.1% offset yield stress was found to reach 38 kg mm -2 with A1-14% AI203 . (430) Creep rates are critically dependent on A1203 concentration with considerable rate reductions being achieved. For example, at 296 K a A1-4 wt % A1203 composite has a creep rate of about 5 x 10 -3 s"1 for the same applied stress. (432) A considerable amount of work has been undertaken on aluminium-fibre composites with boron, graphite, and stainless steel; however, in most cases aluminium alloys rather than high purity aluminium were used as the matrix material. In fibre composites mixture rules relating the elastic modulus of the composite (Ec) and the composite flow stress cr~ to the volume fraction of fibre Vf are

Ec= EfVf + Em ( 1 - Vf)

(15)

with E f the elastic modulus of the fibre and E m the elastic modulhs of the matrix and

o f = o ? (1 - vi)+

vf

(16)

with o ? the flow stress of the matrix at e 1 strain. These very simple relationships have been confirmed for A1stainless steel ffdament composites under tension at room temperature (434) and should provide 'rule of thumb' guide-lines for any Al-flbre composite mixture in which the fibre is considerably stronger than aluminium.

Recovery Work hardened pure aluminium may be easily softened by raising the temperature. Higher temperatures provide the increased thermal energy required to set dislocations in motion. As the temperatures increase, the mobility of the dislocations increase enhancing their chances of annihilation or realignment into lower energy configurations. In this manner, the defect density introduced by a hardening technique at low temperatures (for example, irradiation, cold work, quenching) decreases at higher temperatures to cause reduced mechanical flow stresses and lower electrical resistivity. (The latter is the property measurement by which recovery is usually monitored).

C R Y O G E N I C S . A U G U S T 1972

A number of excellent, well-controlled, irradiationrecovery measurements have been made (435-459) using aluminium, the results of which are summarized in Table 4. Five major stages of recovery may be observed. Although it is not entirely accepted, we have taken the view that the first three stages (I-III) are controlled by I, interstitial and vacancy, in close proximity, annihiliation; II, interstitial migration and subsequent annihilation or recombination; and III, vacancy and vacancy cluster migration. The activation energies have a wide spread within stage III; it is probable that the highest values represent single vacancy motion, while the lower values represent motion energy of either divacancies or of more complex clusters. Recovery experiments (460-471) after cold working at very low temperatures are scarce and are not very thorough. These data are also included in Table 4. If one assumes that all recovery in stages I to III is contributed by interstitials or vacancies or their clusters, one can make an interesting point. [On this assumption, one author (27) thinks that the resistivity decreases are too large to be accounted for solely by point defects or their clusters, and indicates that 'dislocation rearrangements' may contribute.] The results of the experiments beginning with small amounts of deformation at 4 K should be conservative; if one adds an estimate of the vacancies annihilated by interstitial recombination during stages I and II, the ratio would increase to about 4 to 1. Acting to decrease this ratio would be dislocation realignment taking place below room temperature which would contribute to the 40-60% recovery of stage III. A second point to make about cold work introduced at 4 K is that only a small amount of recovery should take place on warming to 76 K, while a much larger amount should occur on warming from 76 K to room temperature. The exact amounts depend, of course, on the degree of initial cold work (higher e promotes earlier dislocation motion), on the purity (higher purities promote dislocation motion at lower temperatures), and on the rate of warming (faster warming rates induce recrystallization at lower temperatures). We have noticed that extensive (> 50%) deformation of 99.999 A1 at 76 K will produce recrystallization at room temperature. Therefore, 99.9999 A1 is expected to begin recrystallization after low temperatur~ deformation somewhere below room temperature. The results of many investigations of high temperature recovery following room temperature deformation have been reported. These are not well-controlled experiments, however, as recovery is continuously taking place during room temperature deformation making it impossible to conduct meaningful isochronal recovery studies. But, such investigations usually reveal that drastic property changes begin to occur about 400 K, depending on degree of deformation, time at temperature, purity, etc. It is probable that future studies will show that stages IV and V converge to the stage III temperature range in very high purity aluminium. Higher temperature recovery experiments using room temperature cold worked specimens have all produced activation energies, (probably relating to dislocation climb) of 1.0 to 1.1 eV. Dislocation annihilation and recrystallization below room temperature, probably reflecting lower activation energies, imply that all is not understood yet concerning the dependence on purity of activation energy.

281

Table 4. Characterization of aluminium recovery I

Defect production technique

References

II

III

Close interstitialvacancy pair Interstitial annihilation migration 10-50 K 50-140 K

IV

Vacancy migration 150-280 K

Electron irradiation

(0.2-3.0 MeV, <10 K)

435-447

40-80% recovered 0.11 eV

5-30% recovered 0.22 eV

Neutron irradiation

(1014-1018 ncm "2, 5 K)

448-459

40-50% recovered

6-25% 28-40% Trapped recovered recovered vacancy 0.30-0.33 eV 0.55-0.63 eV migration, polygonization 280-370 K 3-5% recovery

Deformation

4-8% at 4 K

460, 461

4% recovered

0.5-45% at 78 K

462-465

20-30% recovered 10-15% recovered 0.11-0.32 eV

3-99% at 295 K

466-471

References Reviews

282

10-30% recovered 0.45-0.63 eV range Best value = 0.60 eV Dislocation annihilation, recrystallization 370-600 K 7-20%

40-60% recovered 5-55% Dislocation 370-615 K recovered annihilation, 5-10% 0.39-0.57 eV recrystalliza- recovery tion (99.9999 AI), polygonization (99.995 AI) 25-40% recovery 0.57-0.63 eV 260-380 K Dislocation climb, dislocation annihilation 333--471 K 1.0-1.1 eV

Center, ML-TR-64-280, 1968); previously R. P. Reed, (Cryogenics Div, Boulder, Office of Technical Services, PB171809, 1963) Aerojet Nuclear Systems Co, Materials Properties Data Handbook, Sacramento, Calif (Contract SNP-1, SNPO of NASA and AEC, Cleveland, Ohio, 1970) NATO, Advisory Group for Aeronautical Research and Development Materials Properties Handbook (Technical Dept, Royal Aeronautical Soe, London, 1959) Conrad, H. 'The cryogenic properties of metals', HighStrength Materials (Wiley, New York, 1965) 436 Pearson, T. G., Phillips, H. W. L. 'The production and properties of super-purity aluminium',Met Reviews 2 (1957) 305

The author would like to thank R. P. Mikesell for the assembly and synthesis of some data, Dr H. M. Ledbetter for contributions to the elastic property section, and Dr M. B. Kasen for careful reading of the manuscript. The author is indebted to NASA-Lewis Research Center and the Army, Ft Belvoir for partial financial support of this research.

Fiekett, F. R. 'Aluminium-1. A review of the resistive mechanisms in aluminium',Cryogenics 11 (1971 ) 349 Weiss, V., Sassier, J. G. Aerospace Structural Metals Handbook (Syracuse Univ Press, ASTIA no AD487355, ASD, WPAFB, Ohio, 1963, 1966) Schwattzberg, F. R. Cryogenic Materials Data Handbook (Martin-Marietta Co, Denver, Colo, Technical and Scientific Information Clearinghouse, Defense Documentation

V

Elastic properties 8 9

Goens, E. 'Elastizitatskonstanten des Aluminiumeinkristalls', Ann Phys 17 (1933) 233 Lazarus,D. 'The variation of the adiabatic elastic constants of KC1, NaC1, CuZn, Cu, and A1 with pressure to 10 000 Bars', Phys Rev 76 (1949) 545

CRYOGENICS. AUGUST 1972

10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30

31 32 33 34

35

36 37

Sutton, P. M. 'The variation of the elastic constants of crystalline aluminium with temperature between 63 K and 733 K', Phys Rev 91 (1953) 816 Ramachandran, E. G., Srinivasan, N. 'The determination of the elastic constants of metals by an ultrasonic method', Indian Inst Metals Trans 7 (1953) 173 Long, T. R., Smith, C. S. 'Temperature variation of the adiabatic elastic constants of aluminium', J A coust Soc A m 26 (1954) 146 Schmunk, R. E., Smith, C. S. 'Pressure derivatives of the elastic constants of aluminium and magnesium', J Phys Chem Solids 9 (1959) 100 Kamm, G. N., Alers, G. A. 'Low-temperature elastic moduli of aluminium', J A p p l Phys 34 (1964) 327 Vallin, J., Mongy, M., Salama, K., Beckman, O. 'Elastic constants of aluminium', J A p p l P h y s 35 (1964) 1825 Thomas, J. F. 'The third order elastic constants of aluminium', Phys Rev 175 (1968) 955; also Thesis, Univ Illinois (1968) Ho, P. S., Ruoff, A. L. 'Pressure dependence of the elastic constants for aluminium from 77 to 300 K', J A p p l P h y s 40 (1969) 3151 Huntington, H. B. The Elastic Constants of Crystals, (Academic Press, Solid State Reprints, New York, 1958) Hearmon, R. F. S. 'The elastic constants of anisotropic materials - II',Adv Phys 5 323 (1956) Weertman, J. 'Fast moving edge dislocations on the (111) plane in anisotropic face-centred-cubic crystals', 33 (1962) 1631 Teutonico, L.J. 'The dissociation of dislocations in the face-centered cubic structure', Acta Met 11 (1963) 1283 Sylwestrowicz, W. D., Gibbons, D. F. 'On the temperature dependence of yield stress for aluminium', Phil Mag 3 (1958) 1326 Hill, R. 'The elastic behaviour of a crystalline aggregate', Proc Phys Soc A65 (1952) 349 Wawra, H. H. 'Einfluss der Werkstoffvorbehandlung auf die Elastizit//tseigenschaften vol Aluminiumwerkstoffen', Aluminium 46 (1970) 805 Koster, W. 'Die Temperaturabh~ingigkeit des Elastizit~/tsmoduls reiner Metalle', Z Metallk 39 (1948) 1 Stokes, H.J. 'Apparatus for the measurement of Young's modulus between - 2 0 0 and 700°C by transverse vibration in vacuum', J Sci Instr 37 (1960) 117 Friedel, J., Boulanger, C., Crussard, C. 'Constantes elastiques et froteement interieur de 1'aluminum polygonise', A c t a M e t 3 (1955) 380 Girard, F., Vidal, G. 'Nouvelle m6thode de mesure ~ chaud du coefficient de Poisson des m~taux et alliages', Rev Metall 57 No 2 (1960) 118 Howell, F. M., Stickley, G.W. 'The mechanical properties of ALCOA wrought aluminium alloy products at various temperatures', ASTIA Rept No AD 5100 (1953) 13eniyeva, T. Ya., Larikov, L. N., Poloskiy, I.G. 'Effect of imperfections of crystal structure on elastic and inelastic properties of aluminium', Acad Sci (USSR), Inst Met (1966), Transl FTD-HT-23-784-68, ASTIA Rept No AD 692187 (1969) K/ister, W. 'Elastizita'tsmodul und Dampfung yon Aluminium und Aluminiumlegierungen', Z Metallk 32 (1940) 282 Hordon, M. J., Lement, B. S., Averbach, 13. L. 'Influence of plastic deformation on expansivity and elastic modulus of aluminium', Acta Met 6 (1958) 446 Cook, M., Richards, T. L., 13idmead, G. F. 'Influence of cold deformation on the Young's modulus of some nonferrous metals', J l n s t Metals 83 (1954-55) 41 Wenzl, H., Kerscher, F., Fischer, V., Ehrensperger, K., Papathanassopoulos, K. 'Enfluss yon Neutronenbestrahlung bei 4 K auf den Schermodul yon Kupfer, Aluminium, und Hatin', Z Naturforsh 26A (1971) 489 Ehrensperger, K., Fischer, V., Kerscher, J., Wenzl, H. 'Internal friction maxima due to relaxation of point defects in neutron irradiated aluminium', J Phys Chem Solids 31 (1970) 1835 Lucasson, P. G., Walker, R.M. 'Production and recovery of electron-induced radiation damage in a number of metals', PhysRev 127 (1962) 485 Fine, M.E. 'Apparatus for precise determination of dynamic Young's modulus and internal friction at elevated temperatures', AST1A Rept No AD 126566, Dept Metallurgy, U Northwestern (1957)

C R Y O G E N ICS . A U G U S T 1972

38 39

40

Hughes, D. S., Maurette, C. 'Dynamic elastic moduli of iron, aluminium, and fused quartz', J A p p l Phys 27 (1956) 1184 Voronov, F. F., Vereshchagin, L. F. 'The influence of hydrostatic pressure on the elastic properties of metals. I. Experimental data', Fiz Metal Metalloved 11 No 3 (1961) 443 Folweiler, R. C., Brotzen, F. R. 'The effect of quenched-in vacancies on the elastic modulus of aluminium', Acta Met 7 (1959) 716

Tensile p r o p e r t i e s - single crystals 41 42 43 44 45 46

47 48 49

50

51 52 53 54 55 56 57 58 59 60 61

Taylor, G. L, Elam, C. "The plastic extension and fracture of aluminium crystals', Proc R o y Soc, Set A 108 (1925) 28 Hosford, W. F., Fleiseher, R. L., 13aekofen, W. A. 'Tensile deformation of aluminium single crystals at low temperatures', A c t a M e t 8 (1960) 187 Garstone, J., Honeycombe, R. W. K., Greetham, G. 'Easy glide of cubic metal crystals', Acta Met 4 (1956) 485 Asada, H., Horiuchi, R., Yoshinaga, H., Nakamoto, S. 'Flow stress in A1-Mg alloy single crystals', Trans Inst Metals (Japan) 8 (1967) 159 13enze,F., 13ahler, S. E., Lucke, K. 'Kritische Schubspannung yon Aluminium-Einkris tallen bei h6heren Temperaturer', A c t a M e t I I (1963) 1179 13eruer, R. 'The temperature - and strain rate - dependence of mechanical strength in face-centered-cubic metal single crystals', ASTIA Rept No 284526 (1962); also Z Naturforsch 15a (1960)689 Buck, O., Keefer, D., Robinson, J., Sosin, A., Widersich, H. 'Low temprature defomation of electron irradiated aluminium' A c t a M e t 16 (1968) 195 Deguchi, Y., Yoshida, S. 'Quench hardening in pure aluminium', J Sci Hiroshima Univ, Ser A-ll 30 No 1 - 2 (1966) 21 Dropmann, P., Tensi, H. M., 13orchers, H. 'Kritische Schubspannung und Quergleitspannung von Aluminiumund Aluminium-Magnesium - Einkristallen', Z Metallk 6 l (1970) 848 Grosbras, M., Vergnol, J., Cartraud, M., Villain, J. P., Caisso, J. '~tude des premieres phases de la d~formation de monocristaux d'aluminium falblement alti~', Mere Sci Rev Metallurg 67 (1970) 3 8 7 Haessner, F., Schreiber, D. 'Uber die Verfestigung yon Aluminium - Einkristallen mit geringen Silberzus~/tzen', ZMetallk 48 (1957) 263 Howe, S., Liebmann, 13., L~cke, K. 'High temperature deformation of aluminium single crystals', Acta Met 9 (1961) 625 Jaoul, 13. 'Deformation plastique ~' basse temperature de monocristaux d'aluminium raffine', CR A cad Sci (Paris) 242 (1956) 3039; also J Mech Phys Solids 5 (1957) 95 Jaoul, B., Bricot, L 'Deformation plastique de monocristaux d'aluminium et d'alliages', Mere Sci R ev Metal15 2 (1955) 629 Kelly, A. 'The mechanism of work softening in aluminium', PhilMag l (1956) 835 Kramer, I. R., Podlaseck, S. 'Stress-strain behaviour of aluminium crystals at low pressures', Acta Met l 1 (1963) 7O Lake, J. S. H., Craig, G. 13. 'Substructure strengthening in aluminium', TransASM 61 (1968) 829 Lake, J. S., Craig, G. B. 'Yield of polygonized aluminium single crystals', Met Trans 2 (1971) 1579 Lauriente, M., Pond, R. B. 'Effect of growth imperfections on the strength of aluminium single crystals', J Appl Phys 27 (1956) 950 L/icke, K., Lange, H. 'Uber die Form der Vesfestigungskurve yon Reinstaluminiumkristallen and die Bildung yon Deformationsbandern', Z Metallk 43 (1952) 55 McGrath, J. T., Criag, G. B. 'The effect of striation-type substructure on the deformation of aluminium single crystals', TransAIME 215 (1959) 1022

283

62 63 64 65 66 67 68 69 70 71 72 73 74 75

76

77 78 79 80 81 82 83 84 85 86 87 88 89

284

Miller, R. F., Milligan, W. E. 'Influence of temperature on elastic limit of single crystals of aluminium, silver and zinc', TransAIME 124 (1937) 229 Madden, R., Cottrell, A. H. 'Quench hardening in aluminium single crystals', PhilMag 46 (1955) 735 Mitra, S. K., Dorn, J. E. 'On the nature of strain hardening in face-centred cubic metals', TransAIME 224 (1962) 1062 Noggle, T. S., Koehler, J. S. 'Electron microscopy of aluminium crystals deformed at various temperatures', J Appl Phys 28 (1957) 53 Paxton, H. W., Cottrell, A. H. 'Work hardening in stretched and twisted aluminium crystals', Acta Met 2 (1954) 3 Phillips, W. L. 'Aluminium and copper tested in direct shear', Trans AIME 224 (1962) 845 R6hm, F., Diehl, J. 'Zum Mechanismus der Zugverformung yon EinkristaUen', Z Metallk 43 (1952) 126 Rosenfield, A. R., Averbaeh, B. L. 'Initial stages of plastic deformation in copper and aluminium', Acta Met 8 (1960) 624 Rosi, F. D., Mathewson, C. H. 'A study of the plastic behaviour of high-purity aluminium single crystals at various temperatures', Trans AIME 188 (1950) 1159 Schaefer, R., Nakada, Y., Ramaswami, B. 'The effect of zone refining on stress-strain curves of fee metals', Trans A1ME 230 (1964) 605 Fleischer, R. L., Chaimers, B. 'Size effects in the deformation of aluminium', Trans AIME 212 (1958) 265 Ono, K., Meshii, M. 'Point defect hardening of aluminium: Part I1 Analysis of electron irradiation experiments at 23 K', Trans A S M 60 (1967) 426 Staubwasser, W. 'Uber die Verfestigung yon Aluminium EinkristaUen (99.99% A1) und ihre Deutung', Acta Met 7 (1959) 43 Didenko, D. A., Pustovalov, V. V., Vershinina, V. V. 'Features of the plastic deformation of aluminium single crystals in the temperature range 1.3-4.2 K', Phys Met Metallog (USSR) 23 No 2 (1967) 139 Pustavalov, V. V., Didenko, D. A. 'Singularities of discontinuous flow in aluminium single crystals at low temperatures', Proc Intn Conf on Strength of Metals and Alloys (Tokyo, 1967) 453 Mukherjee, A. K., Mote, J. D., Dorn, J. E. 'Strain hardening of single aluminium crystals during polyslip', Trans AIME 233 (1965) 1559 Sosin, A., Koehler, J. S. 'Electrical resistivity tensor for aluminium single crystals deformed at helium temperature', PhysRev 101 ('1956) 972 Sakui, S., Kakuma, T., Moil, T. 'Behaviour of aluminium single crystals under dynamic loading', Nihon Zinzoku gakkai-shi (Japan Inst Metals) Sendai 29 No 9 (1965) 903 Sato, S., Kelly, A. 'The third stage of work hardening in aluminium crystals deformed at 196 K', Trans AIME 215 (1959) 413 Kramer, 1. R., Demer, L. J. 'The effect of surface removal on the palstic behavior of aluminium single crystals', Trans AIME 221 (1961) 780 Ono, K., Meshii, M., Kauffman, J. W. 'Effect of electron irradiation on mechanical properties of aluminium single crystals at 80 K', A cta Met 12 (1964) 361 Diehl, J., Mader, S., Seeger, A. 'Gleitmechanismus und Oberflffchenerscheinungen bei Kubisch fla'chenzentrierten MetaUen', Z Metallk 46 (1955) 650 Kramer, I. R. 'A study of the effects of surface films on the mechanical properties of metals', WADD Tech Rept 6 0 - 3 1 , ASTIA Rept No AD 246 386 (1960) Nakada, Y., Chalmers, B. 'Effects of surface conditions on the stress-strain curves of aluminium and gold single crystals', TransAIME 230 (1964) 1339 I)iehl, J., Leitz, C., Decker, W. 'Messung der Verfestigung yon Aluminium - Einkristallen bei Tieftemperatur Reaktorbestrahlung', ZMetallk 61 (1970) 443 Howe, S., Elbaum, C. 'The relation between the plastic deformation of aluminium single crystals and polycrystals', PhilMag6 (1961) 37 Kocks, U. F. 'Polyslip in single crystals',Acta Met 8 (1960) 345 Kramer, I. R. 'The effect Of surface removal on the plastic flow characteristics of metals. Part II. Size effects, gold, zinc, and polycrystalline aluminium', TransAIME 227 (1963) 1003

90

Takamura, J. 'Effect of anodic surface films on the plastic deformation of aluminium crystals', Mere Fac Engr Kyoto Univ 18 (1956) 255

Tensile properties o f p o l y c r y s t a l s 91 92 93 94 95

96 97 98 99

100 101 102

103

104 105 106

107 108

109

110

111 112

CottreU, A. H., Stokes, R.J. 'Effects of temperature on the plastic properties of aluminium crystals', Proc R o y Soc (London) A233 (1955) 17 Hirsch, P. B., Warrington, D. H. 'The flow stress of aluminium and copper at high temperatures', Phil Mag 6 (1961) 735 Basinski, Z. S. 'Thermally activated glide in face-centred cubic metals and its application to the theory of strain hardening', PhilMag 4 (1959) 393 Mitra, S. K., Osborne, P. W., Dorn, J. E. 'On the intersection mechanism of plastic deformation in aluminium single crystals', TransAIME 221 (1961) 1206 Nunes, A. C., Rosen, A., Dorn, J. E. 'Effect of strain hardening on the low-temperature thermally activated deformation mechanisms in polycrystalline aluminium', TransASM 58 (1965) 38 Ball, C. J. 'The flow stress of polycrystalline aluminium', PhiIMag 2 (1957) 1011 Bullen, F. P., Rogers, C. B. 'Strain-hardening relationships in polycrystalline aluminium', Phil Mag 11 ( 1965 ) 191 Adler, J. F., Phillips, V. A. 'The effect of strain rate and temperature on the resistance of aluminium, copper, and steel to compression', J l n s t Metals 83 (1954) 80 Arnold, S. V. 'Notch sensitivity and laminated charpy impact strength of 1100-F and 2024-T4 aluminium alloy simulated sheet', Watertown Arsenal Lab, AST1A-AD 225559 (1959) Bell, J. F. 'Single, temperature-dependent stress-strain law for the dynamic plastic deformation of annealed facecentered cubic metals', J A p p l Phys 34 (1963) 134 Carreker, R, P., Hibbard, W. R. 'Tensile deformation of aluminium as a function of temperature, strain rate, and grain size', TransAIME 209 (1957) 1157 Carson, J. M., Hawn, J.M. 'Some properties of iron, copper, and selected aluminium alloys including true stress-true strain at reduced temperatures', A F M L - T R - 6 8 251, Air Force Materials Lab, Wright-Patterson AFB (1968) Cheever, D. L., Howeden, D. G., Monroe, R.E. 'A manufacturing process for producing high-quality electrical strip from ultra high-purity aluminium for magnet applications', A F M L - T R - 6 8 - 3 5 8 , Air Force Materials Lab, Wright-Patterson AFB (1968) Chiddister, J. L., Malvern, L. E. 'Compression-impact testing of aluminium at elevated temperatures', Exp Mech 3 (1963) 81 Derner, P., Kappler, E. 'Uber die Spannungsentfestigung (Work softening-Effekt) an Vielkristallinem Aluminum', . ZNaturforsch 14 No 12 (1959) 1082 Didenko, D. A., Pustovzlov, V. V., Fomenko, V. S. 'Temperature dependence of strength of aiuminium and lead at low temperatures', Phys Chem Mech Mat 5 No 3, (1969) 302 Fleischer, R. L., Hosford, W. F. 'Easy glide and grain boundary effects in polycrystalline aluminium', Trans AIME 221 (1961) 244 Fridlyander, I. N., Ul'yanov, R. A., Nepomnyashchyaya, E. Z., Podkuyko, V. P. 'Mechanical properties of aluminium alloys at low temperatures', Russian Metallurgy No 5 (1967) 72

Fushimi, K., Yonemitsu, H., Okamoto, H., Fukushima, E. 'Tensile properties of various materials at cryogenic temperatures', Tokyo Shibaura Electric Co Ltd, Tokyo, Japan (1970) Hockett, J. E. 'On relating the flow stress of aluminium to strain, strain rate, and temperature', Trans AIME 239 969 (1967); also LA-3544, Los Alamos Scientific Lab, New Mexico (1966) Hansen, N. 'Effect of grain size on the mechanical properties of dispersion-strengthened aluminium, aluminiumoxide products', TransAIME 245 (1969) 1305

Holt, D. L., Babcock, S. G., Green, S. J., Maiden, C. J. 'The strain-rate dependence of the flow stress in some aluminium alloys', Trans A S M 60 (1967) 152

CRYOGENICS. AUGUST 1972

113

114 115 116 117

118 119

120 121 122

123

124 125 126 127 128

129 130 131 132 133 134 135

136 137

Howell, F. M. 'Low temperature properties and applications of aluminium alloys', Rept No 9 - M - 2 1 4 , Alcoa Research Labs, New Kensington, PA (1953) Huitrgten, F. 'Grain boundary and substructure hardening in aluminium', TransAlME 230 (1964) 898 Inglis, N. P., Larke, E. C. 'Strength at elevated temperatures of aluminium and certain aluminium alloys', Proc Inst Mech Engr 172 (1958) 991 Jaoul, B. 'l~tude de la forme des courbes de deformation plastique', JMech Phys Solids 5 (1957) 95 Kadaner, E. S., Kop'ev, I. M., Toropova, L. S. 'Effect of thickness on the mechanical properties of foils made of aluminium and aluminium alloys', Metal Sci and Heat Treatment o f Metals Nos 3 - 4 (1968) 218 Kaufman, J. G., Bogardus, K. O., Wanderer, E.T. 'Tensile properties and notch toughness of aluminium alloys at - 4 5 2 ° F in liquid helium',Adv Cry Engr 13 (1968) 294 Klyavin, O. V., Stepanov, A.V. 'A study of the mechanical properties of solids, in particular of metals at absolute temperatures of 4.2 K and below. 1. Tensile testing of polycrystalline aluminium (99.3%)', Phys Met Metallog (USSR) l (1960) 99 Klyavin, O. V., Stepanov, A. V. 'Effect of surface state on stepwise deformation of aluminium at 1.3 K', Fiz MetalMetalloved 17 No 4 (1964) 592 Kylavin, O. V. 'Effect of deformation rate on the irregular deformation of aluminium at 1.3 K', Phys Met Metallog 17 No 3 (1964) 127 Knoll, H., Maoherauch, E. 'Zum plastischen Verhalten yon vielkristallen Reinstalauminium bei Zugverformung im Temperaturbereich yon 78 his 353 K', Z Metallk 60 (1969) 399 Kochendorfer, A., Wink, W. 'Zugversuche an Stahlen und Nichtersenmetallen bei hohen und tiefen Temperaturen in einer harten Prufmaschine unter Verwendung yon Dehnungs messstreifen zur Kraftme ssung', Archly Eisenhutt 28 (1957 ) 41 Kramer, I. R., Shen, H., Podlaseck, S. E. 'Effect of vacuum on the mechanical behaviour of metals', ASTIA Rept No AD 612022 (1964) Kramer, I. R. 'Effect of specimen diameter on the flow stress of aluminium', TransAIME 239 (1967) 1754 Lindholm, U.S. 'Some experiments with the split Hopkinson pressure bar', JMech Phys Solids 12 (1964) 317 Lindholm, U. S., Yeakley, L.M. 'Dynamic deformation of single and polycrystaUine aluminium', JMech Phys Solids 13 (1965)41 Eindholm, U. S. 'Some experiments in dynamic plasticity under combined stress', Symposium on the Mechanical Behavior of Materials under Dynamic Loads, San Antonio, Texas, US Army Research Office (Durham), Southwest Research Inst; also 'High strain-rate: tension and compression', ExptMech 8 (1968) 1 McAdam, D.J. 'Effect of cold working on endurance and other properties of metals, Part 1', Trans A S S T 8 (1925) 782 Mitra, S. K., Dorn, J. E. 'On the nature of strain hardening in polycrystalline aluminium and aluminium magnesium alloys', TransAIME 227 (1963) 1015 Nock, J. A. 'Properties of commercial wrought alloys', Properties and Physical Metallurgy of Aluminium (1966) North American Aviation Inc. 'Materials property Manual and summary report', AL-2604, AF contract 3 3 ( 6 0 0 ) 28469, Downey, Calif (1957) Pearce, R. "Some aspects of anisotropic plasticity in sheet metals', Int J Mech Sci 10 (1968) 995 Pester, F. 'Festigkeitsprufungen an Stangen und Drahten bei teifen Temperaturen', Z Metallk 24 (1932)67,115 Polosatkin, G. D., Kudryavtseva, L. A., Glazkov, V. M. 'Study of the dynamic yield point of metals at impact velocities up to 1 000 m s"l ',Russian Metallurgy No 5 (1966) 62 Revsin, B., Bodner, S. R. 'Experiments on the dynamic tensile strength of metals and plastics', Technion-Israel Inst Tech, Jaifa, ASTIA Rept No AD 698360 (1969) Rosen, H., Bodner, S. R. 'The influence of strain rate and strain aging on the flow stress of commercially pure aluminium', Technion-lsrael Inst Tech, Haifa, AFOSR 6 6 2587 (ASTIA-AD 641995) (1966)

C R Y O G E N I C S . A U G U S T 1972

138

Shen, H., Podlaseck, S. E., Kramer, 1. R. 'Vacuum effects on the tensile and creep properties of aluminium', Trans AIME 233 (1965) 1933 139 Shiotani, N., Kimura, H., Hasiguti, R. R., Maddin, R. 'The mechanism of hardening in quenched aluminium', A cta Met 15 (1967) 287 140 Sylwestrowicz, W. D., Orowan, E. 'The temperature dependence of the yield stress of copper and aluminium', Mass lnst Tech, Cambridge, Mass, A F O S R - T R - 5 7 - 7 1 , ASTIA-AD 136618 (1957) 141 Trozera, T. A., Sherby, O. D., Dorn, J. E. 'Effect of strain rate and temperature on the plastic deformation of high purity aluminium', TransASM49 (1957) 173 142 Vachet, P., Bonmarin, J. 'Emploi de l'aluminium raffine dans les cryomachines', Rev generale de I'electricite 74 No 6 (1965) 555 143 von Burg, E. 'Mechanisches Festigkeitsverhalten yon Aluminium und seiner Legierungen hei Tieftemperaturen', Schweizer Archly 26 (1960) 110 144 Warren, K. A., Reed, R. P. 'Tensile and impact properties of selected materials from 20 to 300 K', NBS Monograph 63, Supt Documents, US Gov Printing Office, Washington, DC (1963) 145 Wellinger, K., Hofmann, A. 'Prufung metaUischer Werkstoffe in der Kalte', Z Metallk 39 (1948) 233; also Wellinger, K., Seufert, W. 'Untersuchungen uber das Festigkeitsverhalten metaUischer Werkstoffe bei teifen Temperaturen', Z Metallk 41 (1950) 317 146 Westmacott, K.H. 'Hardening in quenched aluminium', PhilMag 14 (1966) 239 147 Westmacott, K.H. 'The flow-stress temperature-dependence of quenched aluminium',Metal Sci J l (1967) 177 148 Yakovleva, E. S. 'The mechanism of plastic deformation and the mechanical properties of aluminium-III. The role of plastic deformation in the formation of the mechanical properties of aluminium', P~z MetalMetalloved 4 (1957) 306 149 Zambrov, J. L., Fontana, M. G. 'Mechanical properties, including fatigue, of aircraft alloys at very low temperatures', TransASM 41 (1949) 480 150 Kochendorfer, A. 'Die Festigkeits - und Form~mderungseigen schafter der Metalle bei teifen Temperaturen', Z Metallk 51 (1960) 73 151 Aust, K. T., Chen, N. K. 'Effect of orientation difference on the plastic deformation of aluminium bycrystals', Acta Met 2 (1954) 632 152 Kato, M., Sogabe, T. 'Grain boundary sliding in highpurity aluminium bicrystal during tensile deformation at high temperature', Trans Jap Inst Metals 12 (1971 ) 107 153 Pond, R. B., Harrison, E. "Grain boundary movement in bicrystalline aluminium', Trans A S M 50 (1958) 994 154 Davis, R. S., Fleischer, R. L., Livingston, J. D., Chalmers, B. 'Effect of orientation on the plastic deformation of aluminium single crystals and bicrystals', TransAIME 209 (1957) 137 155 Fleischer, R. L., Backofen, W. A. 'Effects of grain bounda.ties in tensile deformation at low temperatures', TransAIME 218 (1960) 243 156 Yakovleva, E. S., Yakutorich, M.V. 'The role of grain boundaries in the plastic deformation of aluminium', Dokl A k a d N a u k SSSR 90 (1953) 1027 157 Chin, G. Y., Hosford, W. F., Jr, Backofen, W. A. 'Ductile fracture of aluminium', Trans AIME 230 (1964) 437 158 Armstrong, R., Codd, 1., Douthwaite, R. M., Petch, N. J. 'The plastic deformation of polycrystalline aggregates', PhilMag6 (1961) 45 159 Johnston, T. L., Feltner, C. E. 'Grain size effects in the strain hardening of polycrystals', Met Trans 1 (1970) 1161 160 Kocks, U. F. "The relation between polycrystal deformation and single-crystal deformation',Met Trans 1 (1970) 1121 161 Li, J. C. M., Chou, Y.T. 'The role of dislocations in the flow stress grain size relationships', Met Trans 1 (1970) 1145 162 Conrad, H. 'Work-hardening model for the effect of grain size on the flow stress of metals', Ultrafine Grain Metals (Syracuse Univ Press, Syracuse, NY, 1970) 213 163 Armstrong, R.W. 'Strength properties of ultrafine-grain metals', Ultrafine Grain Metals (Syracuse Univ Press, Syracuse, NY, 1970) 1 164 Yoshida, S., Nagata, N. 'Deformation of aluminium single crystals at high strain rates', Tram Jap Inst Metals 8 (1967) 26 -

285

165 166 167 168 169

170 171 172

173 174

175 176 177

178 179 180 181

182 183

184 185

Kovacs, T. 'Some effects of high rates of strain on the deformation of aluminium single crystals', 2 (1971) 961 Karnes, C. H., Ripperger, E. A. 'Strain rate effects in cold worked high-purity aluminium', JMech Phys Solids 14 (1966) 75 Lindholm, U. S., Yeakley, L. M. 'High strain-rate testing: tension and compression', Expt Mech 8 (1968) 1 McQueen, J. H., Wong, W. A., Jonas, J . J . 'Deformation of aluminium at high temperatures and strain rates', Canad J P h y s 45 (1967) 1225 Nicholas, T., Whitmire, J.N. 'The effects of strain-rate and strain-rate history on the mechanical properties of several metals', ASTIA Rept No 714086, A F M L - T R - 7 0 218 (1970) Dharan, C. K. H., Hauser, F. E. 'Determination of stressstrain characteristics at very high strain rates', Expt Mech 10 (1970) 370 Taborsky, L. 'Some causes of changes in the mechanical properties of aluminium on shock loading', Phys Stat Sol 35 K5 (1969) Bailey, J. A., Singer, A. R. E. 'Effect of strain rate and temperature on the resistance to deformation of aluminium, two aluminium alloys, and lead', J lnst Metals 92 (1963-64) 404 Maraev, S. E., Mudrova, E. I., Elina, N.I. 'Mechanical properties and structure of zone-refined aluminium', Soviet JNon-ferrous Metals 7 No 3 (1966) 87 Borchers, H., Tensi, H. M., Ehrhardt, H. 'Verformungsvermogen von Aluminium und Aluminium-Magnesium Legierungen bei verschiedenen Probenzustanden und Versuchstemperaturen', Aluminium 44 (1968) 546 Wong, W. A. 'Deformation of aluminium at high temperatures and strain-rates', PhD Thesis, McGiU Univ, Montreal (1967) Maeheraueh, E. 'Zur plastischen Verformung yon vielkritsallinem Reinstaluminium', A luminium 37 (1961) 576 Jaoul, B. Clmssard, C. 'Resistance des material - Contribution d l'~tude de la forme des courbes detraction d'eprouvettes monocristallines', CR Acad Sci 234 (1952) 700 Gokyu, K. Kihara, J. 'Work-hardening of aluminium polycrystals', Nippon Kinzoku Gakkai-Shi (Japan lnst Metals J) 31 (1967) 1170 Medrano, R. 'On the temperature dependence of the flowstress in aluminium single crystals', Scripta Met 4 (1970) 315 Derner, P., Kappler, E. 'Zum Temperaturwechselversuch bei der plastischen Verformung der kubisch - fla'chenzentrierten Metalle', Z Phys 163 (1961) 62 Knoll, H., Macherauch, E. 'Temperatur - und Geschwindigtieitspinfluss auf die plastische Verformung yon vielkristallinem Reinstaluminium und einer A1Cu Legierung in Versehiedenen Aushffrtungszusta'nden', Aluminium 45 (1969) 627 Westmaeott, K. H., Smallman, R. E. 'The formation, stability and effect on yield stress of cavities in neutronirradiated aluminium',Mat SciEng 5 (1969/70) 325 Bochirol, L., Btauns, P., Oaudet, G. 'Effects d'irradiations neutroniques ~ 27 K sur les caractdristiques de traction a 27 K sur les caractdristiques de traction a 27 K de l'aluminium', Congres International du Froid, Washington, DC (Sept 1971) Meyers, C. L. 'Strain-hardening effects in BCC, FCC, and HCP metals loaded at very high rates', J M a t Sci 4 (1969) 1 Yoshida, S., Oguchi, A. 'Influence of high hydrostatic pressure on the flow stress of aluminium polycrystals', Trans Japan Inst Metals I 1 (1970) 424

189 190

191 192 193 194 195 196 197 198 199 200 201

202 203 204 205 206 207 208 209 210 211 212 213 214 215

Fatigue properties 186 187 188

286

Moo~e, R. R. 'Some fatigue tests on non-ferrous metals', P r o c A S T M 25 (1925) 66 McCammon, R. D., Rosenberg, H. G. 'The fatigue and ultimate tensile strengths of metals between 4.2 and 293 K ' , P r o c R o y Soc A242 (1957) 203 Awatini, J., Katagiri, K., Koreeda, A. 'Structure changes in metals fatigued at an ultrasonic frequency', Japan Soc Mech EngrBull 12 No 53 (1969) 940

216 217 218

Halford, G. R., Morrow, J. 'Low-cycle fatigue in torsion', ProcASTM 62 (1962) 695 Daniels, N. H. G., Dorn, .I.E. 'The effect of temperature, frequency, and grain size on the fatigue properties of high-purity aluminium', Metals (ASTM Special Technical Publ No 196, Philadelphia, Pa, 1957) p94 Mauer, K. L., Weiss, B. 'Wechselverformung und Ultraschallwechselverformung yon Reinstaluminium und einer AIZnMg-Legierung', A luminium 45 (1969) 609 Thompson, A. W., Baekofen, W. A. 'The effect of grain size on fatigue',Acta Met 19 (1971) 597 Greenail, C. H., Gohn, G. R. 'Fatigue properties of nonferrous sheet metals', Proc A S T M 37 (1937) 160 McAdam, D.J. 'Corrosion-fatigue of non-ferrous metals', P r o c A S T M 27 (1927) 102 Hordon, M.J. 'Fatigue behaviour of aluminium in vacuum', A e t a M e t 14 (1966) 1173 Wright, M. A., Hordon, M. J. 'Temperature dependence of the fatigue transition pressure for aluminium', A cta Met 15 (1967) 430 Hordon, M. J., Wright, M. A. 'Effect of cyclic frequency on the fatigue behaviour of aluminium in vacuum', Tram AIME 242 (1968) 2011 Blucher, J. T., Grant, N. J. 'Low strain rate, high strain fatigue of aluminium as a function of temperature', Trans A1ME 239 (1967) 805 Kennedy, A. J. Processes of Creep and Fatigue in Metals (Wiley, New York, 1963) Gough, H..I., Sopwith, D. C. 'Characteristics of corrosionfatigue of an aluminium specimen consisting of two crystals, J Inst Metals 49 (1932) 93 Wadsworth, N. J. The effect of environment on metal fatigue', Internal Stresses and Fatigue in Metals (Elsevier, New York, 1959) 382; also with Hutchings, J. 'The effect of atmospheric corrosion on metal fatigue', Phil Mag 3 (1958) 1154 Burmeister, R. A., Dodd, R.A. 'The effect of electrodeposited metals on fatigue life', Proc A S T M 62 (1962) 675 Ray, T. K., Seal, A. K. 'Further study of the effect of grain-size on fatigue behaviour of commercially pure aluminium', Trans Indian lnst Metals 20 (1967) 55 Thurston, R.C. 'The mechanism o f fatigue, A review', Trans Canadian MiningMetall 60 (1957) 390 Low, J. R. Jr 'The fracture of metals', Progress in materials science Vol 12 No 1 (Pergamon, New York, 1963) Zackay, V. F., Gerberich, W. W., Parker, E. R. 'Structural modes of fracture', Fracture, Vol 1 (ed H. Liebowitz, Academic Press, New York, 1968) 395 Hancock, J. R., Grosskreutz, J. C. 'Mechanisms of fatigue hardening in copper single crystals', Acta Met 17 (1969) 77 Plumbridge, W. J., Ryder, D. A. 'The influence of specimen geometry on the mode of fatigue crack growth in aluminium aUoys',ActaMet 11 (1969)703 Grosskruetz, J. C.,Wadlow, P. 'Substructure and fatigue fracture in aluminium', A cta Met 11 (1963) 717 McEvily, A. J. Jr, Boettner, R.C. "On fatigue crack propagation in fcc metals',ActaMet 11 (1963) 725 Feltner, C. E. 'Dislocation arrangements in aluminium deformed by repeated tensile stresses', A cta Met 11 (1963) 817 Grosskreutz, J.C. 'Development of substructure in polycrystalline aluminium during constant-strain fatigue', J A p p l Phys 34 (1963) 372 Segall, R. L., Partridge, P. G. 'Dislocation arrangements in aluminium deformed in tension or by fatigue', Phil Mag 4 (1959) 912 Wilson, R. N., Forsyth, P. J. E. 'Some new observations on fatigue damage', JInstMetals 87 (1958-59) 336 Holden, J. 'The formation of sub-grain structure by alternating plastic strain', PhilMag 6 (1961) 547 Holden, J. 'Observations of cyclic structure at large ranges of plastic strain', Acta Met 11 (1963) 691 Snowden, K. U. 'Dislocation arrangements during cyclic hardening and softening in AI crystals', Acta Met 11 (1963) 675 Forsyth, P. J. E., Stubbing-ton, C.A. 'The influence of substructure on the slip observed in pure aluminium and some aluminium alloys when subjected to fatigue stresses', JlnstMetals 83 (1954-55) 173

CRYOGENICS. AUGUST 1972

2t9 220 221 222 223 224 225 226 227 228 229 230 231

232 234 235 236 237 238 239 240 241 242 243 244 245 246 247 248

249

Forsyth, P. J. E. 'Some further observations on the fatigue process in pure aluminium', J l n s t Metals 82 (1953-54) 449 Thomas, K. 'Electron microscope investigation of fatigue', Proc Fifth Conf for Electron Microscopy (Academic Press, New York, 1962) Avery, D. H., Baekofen, W. A. 'Nucleation and growth of fatigue cracks', Fracture of Solids (Vol 20, Met Soc Conf AIME, Interscience, New York, 1963) 339 Smith, G. C. 'The initial fatigue crack', Proc Royal Soc (London) A242 (1957) 189 Alden, T. H., Backofen, W. A. 'The formation of fatigue cracks in aluminium single crystals',Acta Met 9 (1961) 352 Thompson, N. "Some observations on the early stages of fatigue fracture', Fracture (Wiley, New York, 1959) 354 Coffin, L. F., Jr 'Cyclic straining and fatigue', Internal Stresses and Fatigue in Metals (Elsevier, New York, 1959) 363; also 'Low-cycle fatigue', TransASM 55 (1963) 15 Gilman, J. J. Discussion of reference 204, Fracture (Wiley, New York, 1959) 374 Levine, E., Weissmann, S. 'The s~ructure and development of fatigue striations in AI', Trans A S M 61 (1968) 128 Meaking, J. D., Petch, N. J.. 'Atomistic aspects of fracture', Fracture of Solids (AIME Met Soc Conf, Vol 20, Interscience Publ, New York, 1963) 393 Grosskreutz, J. C. 'Residual surface strains in cyclically deformed aluminium and gold',Acta Met 13 (1965) 1269 Terminasov, Y. S., Yar-Mukhamedov, S. K. 'X-ray investigation of fatigue of single crystals of aluminium at normal and low temperatures', ASTIA Rept No 431971 (1963) Glover, H. J., Gordon, S. A., Jackson, L. R. Fatigue of Metals and Structures (Bureau of Aeronautics, Dept Navy, Sutp Documents, US Gov Printing Office, Washington, DC, 1954) Cazaud, R. Fatigue of Metals (Philosophical Lib Publ, New York, 1953) Sines, G., Waisman, J. L. Metal Fatigue (McGraw-Hill, New York, 1959) Parker, E. IL, Fegredo, D.M. 'Nucleation and growth of fatigue cracks', Internal Stresses and Fatigue in Metals (Elsevier, New York, 1959) 263 Wilkov, M. A., Shield, R. 'Mechanism of fatigue in metallic crystals', ASTIA Rept No AD 616341 (1965) Thompson, N. 'Recent work on the nature of fatigue damage', Sym on Fatigue, Czech Sci Techn Soc, Prague (1961) 84 Beachem, C. D. 'Microscopic fracture processes', Fracture Vol 1, (ed H. Liebowitz, Academic Press, New York, 1968) 243 Gough, H. J. Crystalline Structure in Relation to Failure of Metals - Especially by Fatigue (Edgar Marburg Lecture, P r o c A S T M 33 pt II, Philadelphia, Pa, 1933) Weibull, W., Odquist, F. K, G. Colloquium on Fatigue (Springer-Verlag, Berlin, 1956) International Conference on Mechanisms of Fatigue in Solids, Orlando, Florida 1962, Acta Met 11 (1963) 643 Mott, N.F. 'A discussion of work-hardening and fatigue in metals', Proc R o y Soc (London) A242 (1957) 145-227 Raymond, M. H., Coffin, L. F. 'Geometric and hystersis effects in strain-cycled aluminium', Acta Met 11 (1963) 801 Flewitt, P. E. J., tleaid, P.T. 'The growth of fatigue cracks in aluminium', lntn J Fracture Mech 7 (1971 ) 17 Ogura, T., Karashima, S. 'Studies on the subscructure developed around fatigue cracks of aluminium', J Japan InstMetals 34 (1970) 739 Ogura, T., Karashima, S. 'Transmission electron microscopy studies on structure developed around fatigue cracks of aluminium', J Japan Inst Metals 34 (1970) 746 Hori, S., Furushiro, N., Suganuma, T. 'Fatigue of aluminium by uniaxially repeated tensile load', J Jap Inst Light Metals 20 No 9 (1970) 429 Plumbridge, W. J. 'The influence of specimen thickness on fatigue crack propagation in aluminium and aluminium alloys', Quantitative Relation Between Properties and Microstructure (ed D. G. Brandon and A. Rosen, Israel Univ Press, Haifa, 1969) 399 Benham, P. P. 'Fatigue of metals caused by a relatively few cycles of high load or strain amplitude', Metallurgical Reviews, Vol 3, (lnst Metals, London, 1958) 203

C R Y O G E N ICS . A U G U S T

1972

250 251

252 253 254 255 256

257

Zamfik, S. Y., Shewchuk, J. 'Fatigue damage under creep effect', Intn Conf on Pressure Vessel Tech, pt 2, Materials and Fabrication (ASME, New York, 1967) paper 11-89, 1157 Bohmer, M., Munz, D. 'Das Dauerschwingver halten metallischer Wefkstoffe im Vakuum und in verschiedenen Gasatomosph~iren, Teil I',Metall 24 (1970) 446; 24 (1970) 857 Shen, H. 'Effect of vacuum environment on the mechanical behavior of materials', ASTIA Rept No AD 665868, Martin Co, Denver, Colo, MCR67-423 (1967) Wright, M. A., Hordon, M.J. 'Effect of residual gas composition on the fatigue behavior of aluminium', Trans AIME 242 (1968) 713 Borovik,.E.S., Mamedov, M. Sh., Volotskaya, V. V. 'The impact strength of metals', Fiz Metal Metallov 19 No 3 (1965) 451 Borovik, E. S., Mamedov, M. Sh., Volotskaya, V. G. 'Low-temperature strength of metals under pulsed load', NASA Rept No N67-21952 (1966) Ham, J. L. Reiehenback, G. S. 'Fatigue testing of aluminium in vacuum', Symposium on Materials for Aircraft, Missiles, and Space Vehicles, ASTM Spec Tech Publ 345, Philadelphia, Pa (1963) 3 Miller, K. J., Rizk, M. N, 'The interactions of deformation processes in high purity aluminium', Fracture (Chapman and Hall Ltd, London, 1969) 688

Creep p r o p e r t i e s

258 259 260 261

262 263 264 265 266 267 268 269 270 271 272

273

274 275

Weertman, J., Weertman, J. R. 'Mechanical properties, strongly temperature dependent', Physical Metallurgy (Wiley, New York, 1965) 793 Garofalo, E. Fundamentals of Creep and Creep-Rupture in Metals (MacMillan, New York, 1965) Friedel, J. Dislocations (Pergamon, New York, 1964) 303 Bird, J. E., Muckherjee, A. K., Dorn, J. E. 'Correlations between high-temperature creep behavior and structure', Quantitative Relation Between Properties and Microstructure, Proc Intr Conf, Haifa, Israel (Israel Univ Press, Jerusalem, 1969) 255 Lubahn, J. D., Felgar, R. P. Plasticity and Creep of Metals (Wiley, New York, 1963) Creep and Recovery (ASM, Metals Park, Ohio, 1957) Creep and Fracture of Metals at High Temperatures, Symposium held at National Physical Lab (HMSO, London, 1956) Sherby, O. D., Dorn, J. E. 'Creep correlations in alpha solid solutions of aluminium', Trans A1ME 194 (1952) 959 Sherby, O. D., Dorn, J. E. 'Some observations on correlations between the creep behavior and the resulting structures in alpha solid solutions', TransAIME 197 (1953) 324 Sherby, O. D., Orr, R. L., Dorn, J. E. 'Creep correlations of metals at elevated temperatures', Trans AIME 200 (1954) 71 - Ludemann, W. D., Shepard, L. A., Dorn, J. E. 'Recovery of the high-temperature creep properties of polycrystalline aluminium', Trans AIME 218 (1960) 923 Sherby, O. D., Goldberg, A., Dorn, J. E. 'Effect of prestrain histories on the creep and tensile properties of aluminium', T r a n s A S M 4 6 (1954) 681 Huang, H. l-Lieh, Sherby, O. D., Dorn, J. E. 'Activation energy for high temperature creep of high purity aluminium', TransAIME 206 (1956) 1385 Raymond, L., Dorn, J. E. 'Recovery of creep-resistant substructures', Trans AIME 230 (1964) 560 Sherby, O. D., Frenkel, R., Nadeau, J., Dorn, J. E. 'Effect of stress on the creep rates of polycrystalline aluminium alloys under constant structure', TransAIME 200 (1954) 275 Raymond, L., Ludemann, W., Jaffe, N., Dorn, J. 'The effect of decreases in stress on the creep behaviour of polycrystalline aluminium in the dislocation climb region', Trans A S M 54 111 (1961); also WADD Tech Rept 60-326, ASTIA Rept No AD 245 (1960) 309 Sherby, O. D., Lytton, J. L., Dorn, J. E. 'Activation energies for creep of high-purity aluminium', A cta Met 5 (1957) 219 Jaffe, N., Dorn, J. E. 'Effect of stress on the creep rate of high-purity aluminium in the cross-slip region', Trans A1ME 224 (1962) 1167

287

276 277

278 279 280

281 282 283 284 285 286 287 288 289 290 291 292 293 294

295 296 297 298 299 300 301 302 303 304 305

288

Harper, J., Dorn, J. E. 'Viscous creep of aluminium near its melting temperature', Acta Met 5 (1957) 654 Dorn, J. E. 'Some fundamental experiments on high temperature creep', JMech Phys Solids 8 (1954) 85 ; also Creep and Fracture of Metals at High Temperatures (1954 Symposium at Nat Phy Lab, HMSO, London, 1956) 89 Servi, I. S., Grant, N. J. 'Creep and stress rupture behavior of aluminium as a function of purity', TransAIME 191 (1951) 909 Jenkins, W. D. 'Creep of high-purity aluminium', J Res NBS 46 (1951) 310 Yampolski, B. Ya., Amflteatrova, T. A. 'Investigation of the deformation of metals under small stresses on some regularities of creep in copper and aluminium', Phys Met Metallog4 No 1 (1957) 106 Simmons, ff~ F., Crossd H.C. 'Creep of 2S-0 aluminium sheet at 400 and 450 C', Battelle Memorial lnst Rept B M I - T - 2 9 (June, 1950) Sweetland, E. D., Parker, E. R. 'Effect of surface condition on creep of some commercial metals', J A p p l Mech 20 (1953) 30 Betteridge, W. 'The low-stress torsional creep properties of pure aluminium', J l n s t Metals 82 (1953-54) 149 Weertman, J. 'Creep of polycrystalline aluminium as determined from strain rate tests', JMech Phys Solids 4 (1956) 230 Wyatt, O.H. 'Transient creep in pure metals',Proc Phys Soc B66 (1953) 495 Ahlquist, C. N., Nix, W. D. 'The measurement of internal stresses during creep of AI and AI-Mg alloys', Acta Met 19 (1971) 373 McLean, D. 'Crystal slip in aluminium during creep', J l n s t Metals 81 (1952-53) 133 McLean, D. 'Grain-boundary slip during creep of aluminium', J1nstMetals 81 (1952-53) 293 Dushman, S., Dunbar, L. W., Huthsteiner, H. 'Creep of metals', J A p p l Phys 15 (1944) 108 Dorn, J. 'The spectrum of activation energies for creep', Creep and Recovery (ASM, Cleveland, Ohio, 1957) 255 Fleischer, R. L., Backofen, W. F. 'Effect of orientation on creep of aluminium single crystals at 4.2 K', TransAIME 218 (1960) 188 Koval, V. A., Soldatov, V. P., Startsev, V. 1. 'Creep of aluminium in the temperature range 1.4-4.2 K', Soviet Phys Solid St 12 No 10 (1971) 2347 Schwope, A. D., Shober, F. R., Jackson, L. R. 'Creep in metals', Nat Adv Comm for Aeronautics, Tech Note 2618 (1952) Johnson, R. D., Young, A. P., Schwope, A. D. 'Plastic deformation of aluminium single crystals at elevated temperatures', Nat Adv Comm for Aeronautics Tech Note 3351 (1955). Miehelitsch, M. 'Uber das Kreichen kubisch fla'chenzentrierter MetaU-Einkristalle', Z Metallk 50 (1959) 548 Weertman, J. 'Creep of aluminium single crystals', J A p p l Phys 27 (1956) 832 Underwood, E. E., Marsh, L. L. 'Effects of alloying elements on plastic deformation in aluminium single crystals', TransAIME 206 (1956) 477 Garofalo, F. 'An empirical relation defining the stress dependence of minimum creep rate in metals', Trans AIME 227 (1963) 351 Cottrell, A. H. 'The time laws of creep', JMech Phys Solids 1-2 (1952-4) 53 Servi, I. S., Grant, N. S. 'Structure observations of aluminium deformed in creep at elevated temperatures', Trans AIME 191 (1951) 917 McLean, D. 'Crystal fragmentation in aluminium during creep', JInstMetals 81 (1952-53) 287 McLean, D. 'Creep processes in coarse-grained aluminium', JlnstMetals 80 (1951-52) 507 Gervais, A. M., Norton, J. T., Grant, N. J. 'Subgrain formation in high-purity aluminium during creep at high temperatures', Trans AIME 197 (1953) 1166 McLean, D. 'Interaction between crystal slip and grain boundary sliding during creep', Creep and Fracture of Metals at High Temperature (HMSO, London, 1956) 73 Wilms, G. R. 'Some observations on the creep of prestrained aluminium', J Inst Metals 83 (1954-55) 427

306 307 308 309 310 311

Wood, W. A., Rachinger, W. A. 'The mechanism of deformation in metals, with special reference to creep', J Inst Metals 76 (1949-50) 237 Wood, W. A., Suiter, J. W. 'Stress-recovery in aluminium', J l n s t Metals 80 (1951-52) 501 Servi, 1. S., Norton, J. T., Grant, N.J. 'Some observations of subgrain formation during creep in high purity aluminium', TransAIME 194 (1952)965 Stein, C. 'The effect of quenching, irradiation damage, and prior fatigue on the creep of pure aluminium', Trans AIME 245 (1969) 2251 Chang, H. C., Grant, N.J. 'Inhomogeneity in creep deformation of coarse grained high purity aluminium', Trans A1ME 197 (1953) 1175 Brunner, H., Grant, N. J. 'Measurement of deformation resulting from grain boundary sliding in aluminium and alummmm-magnesmm from 410 F to 940 F , TransAIME 218 (1960) 122 Ahlquist, C. N., Gasca-Neri, R., Nix, W. D. 'A phenomenlogical theory of steady state creep based on average internal and effective stresses', A c t a M e t 18 (1970)663 Ahlquist, C. N., Nix, W. D. 'The measurement of internal Stresses during creep of A1 and A1-Mg alloys', Acta Met 19 (1971) 373 Nix, W. D., Barrett, C. R. 'Interpretation of steady state creep in terms of the relation between dislocation glide and recovery', Trans A S M 61 (1968) 695 Gasea-Neri, R., Ahlquist, C. N., Nix, W. D. 'A phenomenlogical theory of transient creep', Acta Met 18 (1970) 655 Rosen, A. 'Delay time during creep of aluminium', J l n s t Metals 99 (1971) 111 Schoeck, G. 'Theory of creep', Creep and Recovery (ASM, Metals Park, Ohio, 1957) 199 Lytton, J. L., Shepard, L. A., Dorn, J. E. 'The activation energies for creep of single aluminium crystals favorably oriented for (111) [1011 slip', TransAIME 209 (1958) 220 Schwope, A. D., Jackson, L. R. 'A survey of creep in metals', NACA Tech Note 2516, Washington, DC (1951) Dorn, J. E. 'A survey of recent results on experimental determinations of activation energies for creep', Fourth Sagamore Ordnance Mat Res Conf, Racquette Lake, NY (1957) 31 Stein, C. 'The role of subgralns during the creep of pure aluminium', Sandia Lab Rept S C - R R - 6 7 - 2 9 7 5 , Albuquerque, New Mexico (1967) Sellars, C. M., Tegart, W. J. MeG. 'La relation entre la rdsistance et la structure dans la ddformation a' chaud', Mem Sci Rev MetaUurg 63 731 (1966) Voloshina, L. A., Rozenberg, V. M. 'Investigation of creep in bicrystals of aluminium', Fiz Metal Metalloved 12 No 1 (1961) 118 Voloshina, L. A., Rozenberg, V. M. 'Effect of orientation of aluminium single crystals on creep', Fiz Met Metalloved 13 No 3 (1962) 474 Van Echo, J. A., Simmons, W. F., Cross, H. C. 'Creep of 2S-0 aluminium sheet at 500 and 550 C', Battelle Mem Inst Rept No BM1-766 (Sept 1952) McLean, D., Farmer, M. H. 'The relation during creep between grain-boundary sliding, sub-crystal size, and extension', J l n s t M e t 85 (1956-57) 41 Guard, R. W., Hibbard, W. R. 'Tensile creep of high purity aluminium', Trans AIME 206 (1956) 195 Rocher, Y. A., Shepard, L. A., Dorn, J. E. 'Activation energies for creep of single aluminium crystals favorably oriented for cubic slip', TransAIME 215 316 (1959) •

312 313 314 315 316 317 318 319 320

321 322 323 324 325 326 327 328

.

-

O

O

,

Hardness

329 330 331 332

Guiilet, L., Cournot, J. 'Sur la variation des proprietes mdcaniques de quelques mdtaux et alliages aux basses tempdratures', Rev Metal119 (1922) 215 Petty, E. R. 'A low-temperature inflection in the temperaturedependence of hardness of pure metals', J Inst Metals 89 (1960) 123 Sauerwaid, F. 'Die Abhangigkeit den Hffrte yon der Temperatur', ZMetallk 45 (1924) 315 Kurth, A. 'Untersuchungen//ber des Einfluss der W~rrne aud die Hffrte der Metalle', Z Vereines deutlngenieure 53 (1909) 209

CRYOGENICS

. AUGUST

1972

333

334 335 336 337 338

339 340 341 342

Tweer-Schriner, i., Lintner, K. 'Bestrahlungsinduzierte H~te~'aderung verschieden verformter Aluminium-Proben', Akad ICissenschaflen, Vienna. Math-Natur Klasse, Anzeiger 15 No 105 (1968);MetaleAbs (1969) 1498 Kasen, M, B., Reed, R. P. 'Ahaminium hardness as a function of purity', unpublished data (1971 ) Helling, W., Neunzig, H. 'Development and possible uses of high purity Al',Metall 5 (1951)424 Tomlinson, J. E. Unpublished data, reported in reference 6 (1957) Bettler, H. 'Raffinal - its preparation and application', Aluminium Swisse 3 (1953) 51 Kimura, H., Nakano, O. "Effects of addition of V a and Via Group elements on recrystallization and mechanical properties of high purity aluminium', Trans Nat Res Inst for Metals 13 No 2 (1971) 13 lvafi Ko, A. A. 'Handbook of hardness data', Acad Sci Ukrainian SSR, Transl distribuyed by Nat Tech Inform Ser, Springfield, Va. Rept No TT70-50177 (1968) Upit, G. P., Varehenya, S. A. 'Hardness of mono- and polycrystalline aluminium with loads of 0.3 to 10 000 grammes', Russian Metallurgy No 2 (1969) 107 Hikage, T. 'On the microhardness of aluminium single crystals', Tohoku Univ Sci Repts, Set A 5 (Sendai, Japan, 1953) 254 Cotner, J. R., Tegart, W. J. MeG. 'High-temperature deformation of aluminium-magnesium alloys at high strain rates', JlnstMetals 97 (1969) 73

357

358 359 360 361 362

Strengthening-work hardening 363 364 365

Impact 343

344 345

Fontana, M.G. 'Mechanical properties and physical metallurgy of aircraft alloys at very low temperatures, Pt I', ASTIA Rept No 24533, USAF Tech Rept No 5662 (1948) Johnston, H. L., Brooks, H. E. 'Impact strength of various metals at temperatures down to 20 absolute', ASTIA Rept No 67, Ohio State Res Round (1952) Coibeek, E. W., MaeGillivray, W. E. 'The mechanical properties of metals at low temperatures, Pt II - Non-ferrous materials', Translnst Chem Engr (London) I l (1933) 107

Strengthening-alloying 346

347 348 349 350

351 352 353

354 355 356

Nix, W. D., Menezes, R. A. 'Physics of strengthening mechanisms in crystalline solids', Annual Review of Materials Science, Vol 1 (Annual Reviews, Inc, Palo Alto, Calif, 1971) 313 Cottrell, A. H. 'Interactions of dislocations and solute atoms', Relation of Properties to Microstructure (ASM, Metals Park, Ohio, 1954) 131 Fleiseher, ILL. 'Solid solution hardening', The Strengthenhag of Metals (Reinhold Publ Col, New York, 1964) 93 Flinn, P. A. 'Solid solution strengthening', Strengthening Mechanisms ha Solids (ASM, Metals Park, Ohio, 1962) 17 Haasen, P. 'Structural defects in solid solutions', Alloying Behaviour and Effects in Concentrated Solid Solutions (Met Soc Conf, AIME, Vol 29, Gordon and Breach, New York, 1965) 270 Haasen, P. 'Mechanical properties of solid solutions and intermetallic compounds', Physical Metallurgy (ed R. W. Cahn, Wiley, New York, 1965) 821 Fine, M. 'Precipitation hardening', The Strengthening of Metals (Reinhold Publ Co, New York, 1964) 141 Hart, E. W. 'Theories of dispersion hardening', Relation of Properties to Microstructute (ASM, Metals Park, Ohio, 1954) 95; see also, Dorn, J. E., Start, C. D. Effect of Dispersions on Mechanical Properties, 71 Kelly, A. 'Theories of precipitation and dispersion strengthening', Electron Microscopy and Strength of Crystals (Wiley, New York, 1963) 947 Tensi, H. M., Dropmann, D., Botchers, H. 'Plastisches Verhalten yon Aluminium-Magnesium-Einkristallen', ZMetalik 61 (1970) 518 Dora, J. E., Pietrokowsky, P., Tietz, T. E. 'The effect of alloying elements on the plastic properties of aluminium alloys', TransAIME 188 (1950) 933

CRYOGENICS. AUGUST 1972

Sherby, O. D., Anderson, IL A., Dotn, J. E. 'Effect of alloying element on the elevated temperature plastic properties of alpha solid solutions of aluminium', Trans AIME 191 (1951) 643 Krupotkin, Ya. M. 'Effect of cerium on the mechanical properties and electrical conductivity of Al', Energetika 12 (1968) 111 Krupotkin, Ya. M., Botts, S. M. 'The mechanical properties of AV000 aluminium with small additions of different elements', Metal Sci Ht Treatment 7 - 8 (1969) 632 Bryne, J. G., Fine, M. E., Kelly, A. 'Precipitate hardening in an aluminium-copper alloy', Phil Mag 6 (1961 ) 1119 Dash, J., Fine, M. E. 'Temperature and composition dependence of the strength of aluminium base fine alloy single crystals',Acta Met 9 (1961) 149 Dew-Hughes, D., Robertson, W. D. 'The mechanism of hardening in aged aluminium-copper alloys', Acta Met 8 (1960) 156

366 367 368 369 370 371 372 373

Berner, IL, Kronmuller, H. 'Plastische Verformung yon Einkristallen', Modern Probleme der Metallphysik (SpringerVerlag, Berlin, 1965) 35 ICronmuller, H. 'Theorie der plastischen Verformung', Moderne Probleme der Metallphysik (Springer-Verlag, Berlin, 1965) 126 Parker, E. R. 'Strain hardening', The Strengthening of Metals (Reinhold, New York, 1964) 77 Work Hardening, Met Soc Conf, Vo146 (ed 1. P. Hirth, J. Weetman, Gordon and Breach, New York, 1966) Proceedings of the International Conference on the Strength of Metals and Alloys, Tokyo, Japan, 1967 (Suppl to Trans Japan Inst Metals 9 1968) Second International Conference on the Strength of Metals and Alloys (Proc Am Soc Metals Conf, 1970) Betekhfin, V. J., Petrov, A. I. 'Mosaic block misorientation and the strengthening of aluminium by low-temperature predeformation', Fiz MetalMetaUoved 25 No 4 (1968) 726 Nost, B , Nes, E. 'Dislocations in aluminium under stress observed by Lang x-ray topography', Acta Met 17 (1969) 13 Rider, J. G., Foxon, C. T.B. 'An experimental determination of electrical resistivity of dislocations in aluminium', Phil Mag 13 (1966) 289 Hammad, F. H., Nix, W. D. 'Dislocation configurations developed in aluminium during high-temperature deformation', Trans ASM 59 (1966) 94 Hoidon, M. J., Averbaeh, B. L. 'X-ray measurements of dislocation density in deformed copper and atuminium single crystals',Aeta Met 9 (1961) 237

Strengthening-substructure 374 375 376

377 378

Kelly, A. 'An x-ray microbeam study of polycrystalline specimens of aluminium and iron deformed in tension', Acta Cryst 7 (1954) 554 Swann, P. R. 'Dislocation arrangement in face-centered cubic metals and alloys', Electron Microscopy and Strength of Crystals (Interscience, New York 1963 ) 131 Hockett, J. E., McQueen, H. J. 'The effects of strain rate and temperature upon the microstructure and strength of aluminium', Paper 15 A. 3, Second Inter Conf on the Strength of Metals and Alloys (ASM, Metals Park, Ohio, 1970) 991 Wong, W. A., MeQueen, IL J., Jonas, J. J. 'Recovery and recrystallization of aluminium during extrusion', J Inst Metals 95 (1967) 129 Alterpohl, D. 'Neuere Untersuchungen uber Subkorner in Aluminium', ZMetallk 49 (1958) 331

Strengthening-irradiation 379 380

Kelly, B.T. Irradiation Damage to Solids (Pergamon, New York, 1966) Holmes, D. IC 'Radiation damage in non-fissionable metals', The Interaction of Radiation with Solids (Proc Intern Summer School on Solid State Plays, Mol, Belgium, 1963, Interscience, New York, 1964) 147

289

381

382 383 384 385 386

Wechsler, M. S. 'Radiation embrittlement of metals and alloys', The Interaction of Radiation with Solids (Proc Intern Summer School on Solid State Phys, Mol, Belgium, 1963, lnterscience, New York, 1964) 296 Damask, A. C., Dienes, G.J. Point Defects in Metals (Gordon and Breach, New York, 1963) Radiation Damage in Solids (Proc lntn School of Physics, Enrico Fermi, Course XVIII, Academic Press, New York, 1962) International Conference on Vacancies and Interstitials in Metals, NASA N69-12994-13028, Clearinghouse, Springfield, Va (1968) Ono, K., Meshii, M. 'Work-hardening of irradiated aluminium', Work Hardening, Met Soc Conf, Vo146 (Gordon and Breach, New York, 1966) 226 Keefer, D., Sosin, A. 'Electron irradiation hardening of aluminium', Acta Met 12 (1964) 969

408

409 410 411 412 413

Strengthening-quenching 387 388 389 390 391 392 393 394

Altenpohl, D. 'Leerstellen in Aluminium', Aluminium 37 (1961) 401 Hirsch, P. B., Silcox, J., Smallman, R. E., Westmaeott, K. H. 'Dislocation loops in quenched aluminium', Phil Mag 3 (1958) 897 Smallman, R. E., Westmacott, K. H., Coiley, J. C. 'Clustered vacancy defects in some face-centered cubic metals and alloys', JhtstMet 88 (1959-60) 127 Kuhlmann-Wilsdorf, D., Wilsdorf, H. G. F. 'On the behavior of thermal vacancies in pure aluminium', JAppl Phys 31 (1960) 516 Vandervoort, R., Washburn, J. 'On the stability of the dislocation substructure in quenched aluminium', Phil Mag 5 (1960) 24 Loretto, M. H., Ciareborough, L. M., Humble, P. 'The nature of dislocation loops in quenched aluminium', Phil Mag 13 (1966) 953 Edington, J. W., Smallman, R. E. 'Faulted dislocation loops in quenched aluminium',PhilMag 11 (1965) 1109 Cotterill, R. M. J., Segall, R. L. 'The effect of quenching history, quenching temperature and trace impurities and vacancy clusters in aluminium and gold', Phil Mag 8 (1963)

414 415 416 417 418

Strengthening-composites 419 420

1105

395 396 397 398 399 400 401 402 403 404 405 406 407

290

Tartour, J-P, Washburn, J. 'Climb kinetics of dislocation loops in aluminium', PhilMag 18 (1968) 1257 Dobson, P. S., Goodhew, P. J., Smallman, R. E. 'Climb kinetics of dislocation loops in aluminium', Phil Mag 16 (1967) 9 Yoshida, S., Shimomura, Y., Kiritani, M. 'Dislocation loops containing a stacking fault in quenched super-pure aluminium', JPhys Soc (Japan) 17 (1962) 1196 Yoshida, S., Shimomura, Y. 'New type dislocation loops in quenched aluminium',JPhysSoc (Japan) 18 No 11 (1963) 1590 Edington, J. W., West, D. R. 'Two-triangle loops in quenched aluminium', JAppl Phys 37 (1966) 3904 Edington, J. W., West, D. R. 'Three-layer defects in quenched aluminium', Phil Mag 14 (1966) 603 Westmaeott, K. H., Smallman, R. E., Dobson, P. S. 'The annealing of voids in quenched aluminium and a determination of the surface energy', Metal Science J 2 (1968) 177 Edington, J. W., West, D. R. 'Four-layer defects in quenched aluminium', PhiI Mag 15 (1967) 299 Yoshida, S., Kiritani, M., Shimomura, Y. 'Dislocation loops with stacking fault in quenched aluminium', JPhys Soc (Japan) 18 No 2 (1963) 175 Kiritani, M., Yoshida, S. 'A new type aggregation of quenchedin vacancies in aluminium', JPhys Soc (Japan) 18 (1963) 915 Kiritani, M. 'Formation of voids and dislocation loops in quenched aluminium', JPhys Soc (Japan) 19 No 5 (1964) 618 Kiritani, M., Shimomura, Y., Yoshida, S, 'Shape of voids in quenched aluminium', JPhys Soc (Japan) 19 No 9 (1964) 1624 Das, J., Washburn, S. 'Defects formed from excess vacancies in aluminium', Phil Mag (1965) 955

Strudel, J. L. B., Washburn, J. 'The formation of voids and dislocation loops in quenched aluminium single crystals', The Interaction Between Dislocations and Point Defects 1I, report of conference held at HarweU, England (1968) 367; also 'Dislocation substructure in quenched aluminium single crystals', Masters Thesis, U of California, Berkeley (1963) DiMelfi, R. J., Siegel, R.W. 'Effect of impurities upon the nucleation of dislocation loops in quenched aluminium', PhilMag 24 (1.971) 279 Shiotani, N., Kimura, H., Hasiguti, R. R., Maddin, R. 'The mechanism of hardening in quenched aluminium', Acta Met 15 (1967) 287 Tanner, L, E. 'On the temperature dependence of flow stress in quenched aluminium crystals', Acta Met 8 (1960) 730 Birnbaum, H. K. 'On the temperature dependence of the flow stress in quenched aluminium crystals', Acta Met 9 (1961) 968 Yoshida, S., Mumkami, H. 'Annihiliation of faulted dislocation loops in quenched aluminium', Tokyo Univ, Fac of Engr J 43 (1968) 54 Kiritani, M. 'Nucleation and growth of secondary defects in quenched face-centered cubic metals', JPhys Soc Japan 20 (1965) 1834 Guimbard, J., Descouts, P., Gobija, P. 'The influence of ageing on the hardening of pure aluminium quenched to low temperature', Brit JAppl Phys 17 (1966) 1433 Guimbard, J., Gobin, P. 'The influence of quenching temperature on the hardening of pure aluminium quenched to -72°C ', Brit JAppl Phys 18 (1967) 1641 Tanner, L. E., Maddin, R. 'Overshooting in quenched aluminium crystals', Acta Met 7 (1959) 76 Mori, T., Meshii, M. 'Plastic deformation of quenchhardened aluminium single crystals', Acta Met 17 (1969) 167

421 422

423 424 425 426 427 428 429 430 431 432 433

Davis, L. W., Bmdstteet, S.W. Metal and Ceramic Matrix Composites (Cahners, Boston, Mass, 1970) Structural design guide for advanced composite apphctions', prepared by North American Rockwell Corp, Los Angeles, Calif, limited distribution by Air Force Materials Lab, Wright-Patterson AFB, Ohio (1971) Sinclair, P. M. 'Composites: designers wait and contemplate', Industrial Research (Oct 1969) 59 Hanby, K. R. 'Fiber-reinforced metal-matrix composites', DMIC Repts No 241, S-21, S-27, Defense Metals - Inform Center, Battelle Mem Inst, Columbus, Ohio (1967, 1968, 1969) Sherby, O. D., Shyne, J. C. 'Research on metallic base composites',NavalResearch Reviews 23 No 5 (1970) 12 Grant, N. J. 'Dispersed phase strengthening', The Strengthening of Metals (Reinhold, New York, 1964) 163 Tegart, W. J. MeG. 'The effect of second phases on the mechanical properties of metals and alloys', J A ust Inst Metals 15 (1970) 47 'Metal-matrix composites', A1ME Symposium, DMIC Memo 243, Defense Metals Inform Center, Battelle Mere Inst, Columbus, Ohio (1969) ltansen, N. 'Dispersion strengthening of aluminium aluminium - oxide products', Acta Met 18 (1970) 137 ttansen, N. 'Oxide dispersion strengthening in sintered aluminium products, SAP', Met Trans 1 (1970) 545 Hansen, N. 'Strengthening of aluminium by a three-dimensional network of aluminium-oxide particles', A cta Met 17 (1969) 637 Kylavin, O.V. 'Strength properties of material type SAP at 4.2-1.3 K', Fiz MetalMetalloved 19 No 6 (1965) 937 Guyot, P., Debeir, R. 'Durcissement a basses temperatures dans les alliages du type SAP/AI-A1203', Acta Met 14 (1966) 43 Milicka, K., Cadek, J., Rys, P. 'Creep of aluminium strengthened by alumina particles', Acta Met 18 (1970) 733 Ansell, G. S., Kim, H. S. 'Fracture behaviour of a dispersion strengthened SAP-type alloy', AST1A Rept No AD 282 509 (1962)

C R Y O G E N I C S . A U G U S T 1972

434

Pinnel, M. P., Lawley, A. 'Correlation of uniaxial yielding and substructure in aluminium-stainless steel composites', Met Trans 1 (1970) 1337

452 453

Recovery 435 436 437 438 439 440

441 442 443 444 445 446

447

448

449

450 451

Chaplin, R. L., Simpson, H. M. 'Energy dependence of the stage I recovery of aluminium', Phys Rev 163 (1967) 587 Simpson, H. M., Chaplin, R. L. 'Interstitial migration in aluminium', Phys Rev 187 (1969) 858 Sosin, A., Garr, K. R. Recovery of electron-irradiated aluminium and atuminium alloys, stage I', Phys Rev 161 (1967) 664; stage I1, Phys Rev 162 (1967) 681 Sosin, A., Rachal, L. H. 'Recovery study in pure and alloyed aluminium following electron irradiation', Phys Rev 130 (1963) 2238 Wagner, H., Dworschak, F., Schilling, W. 'Comparison of lattice-parameter and resistivity change in electron-irradiated aluminium', Phys Rev B2 (1970) 3856 Dawson, H. 1., Iseler, G. W., Kauffman, J. W. 'Low temperature recovery studies in electron irradiated aluminium', Lattice Defects and Their Interactions (Gordon and Breach, New York, 1967) 681 Longshore, R. E., Chaplin, R. L. 'Electron irradiation of single crystals of aluminium', J Appl Phys 40 (1969) 351 Peters, P. B., Shearin, P. E. 'Recovery of resistivity of pure and alloyed aluminium in stages II and Ili after 2-MeV electron irradiation', PhysRev 174 (1968) 691 Keefer, D. W°, Robinson, J. C., Sosin, A. 'Modulus effects in metals after low temperature electron irradiation III. Aluminium', Acta Met 16 (1968) 927 Budin, C., Lucasson, P. 'l~tude par trempe et par irradiation de defauts reticulaires dans l'aluminium', Phys Lett 16 (1965) 229 Lwin, Y. N., Doyama, M., Koehler, J. S. 'Stage III annealing study of electron-irradiated pure aluminium', Phys Rev 165 (1968) 787 Doyama, M., Koehler, J. B., Lwin, Y. N., Ryan, E. A., Shaw, D.G. 'Annealing study of dilute aluminium alloys electron irradiated at liquid nitrogen temperature', Phys Rev B3 (1971) 1069 Burger, G., lsebeck, K., V(Jikl, J., Schilling, W., Wenzl, H. 'Erholung des elektrischen Widerstandes reiner Metalle nach Neutronenbestrahlung mit niedriger Dosis bei 4.6 K im Reaktor', Z Angew Phys 22 (1966-7) 452 Dimitrov-Frois, C., Dimitrov, O. 'Influence de quelques dl~ments d'addition sur la restauration de l'aluminium irradid aux neutrons a 78 K', Mem Sci Rev Metall 65 (1968) 425 Sepp, A., Himmler, U., Peisl, H., Wenzl, H., Waidelieh, W. 'Thermal annealing of lattice parameter change in lowtemperature neutron-irradiated aluminium', Phil Mag 14 (1966) 131 lsebeck, K., Muller, R., Schilling, W., Wenzl, H. 'Reaction kinetics of stage III recovery in aluminium after neutron irradiation', Phys Star Sol 18 (1966) 427 Isebeck, K., Rau, F., Schilling, W., Sonnenberg, K., Tischer, P., Wenzl, H. 'Stored energy, volume, and resistivity change in neutron irradiated aluminium', Phys Stat Sol 17 (1966) 259

CRYOGENICS. AUGUST 1972

454 455 456 457 458 459 460 461

462 463 464 465

466 467 468 469 470 471

Ceresara, S., Elkholy, H., Federighi, T., Pieragostini, F. 'On the sub-structure of stage l n - 3 in neutron irradiated AI', Phys Lett 16 (1965) 8 Federighi, T., Ceresara, S., Pieragostini, F. 'On the recovery of aluminium after neutron irradiation at 78 K', PhilMag 12 (1965) 1093 Frois, C., Dimitrov, O. 'Elimination des defauts dans l'aluminium irradie par les neutrons rapides dans l'azote liquide', Mem Sci Rev Met 62 (1965) 697 Burger, G., Meissner, H., Schilling, W. 'The influence of the initial defect concentration on the annealing of lowtemperature irradiated metals', Phys Star Sol 4 (1964) 267 Swanson, M. L., Piercy, G. R. 'Low-temperature neutron irradiation damage in aluminium and gold', Can J Phys 42 (2) (1964) 1605 Ceresara, S., Federighi, T., Gelli, D., Pieragostini, F. 'Recovery of aluminium after neutron-irradiation at 78 K', Nuovo Omento 29 (1963) 1244 McReynolds, A. W., Augustyniak, W., MeKeown, M., Rosenblatt, D. B. 'Neutron irradiation effects in Cu and AI at 80 K', Phys Re~ 98 (1955) 418 Swanson, M. L. 'Low-temperature recovery of deformed aluminium', Can JPhys 42 (1964) 1890 Takamura, S., Okuda, S. 'Electrical resistivity measurements of fcc metals deformed at 4.2 K', J Phys Soc Japan 25 (1968) 714 Lugscheider, W., Breitenfellner, F., Lihl, F. "Erholung yon Gitterdefekten in Aluminium und Legierungen des Aluminiums mit Beryllium, Nickel, und Titan nach einer Zugverformung bei 78 K', Z Metallk 59 (1968) 22 Ceresara, S., Elkholy, H., Federighi, T. 'Influence of the amount of strain at 78 K on the recovery process in A1 99.995%', PhilMag 12 (1965) 1105 Frois, C. 'Elimination des defants crees par laminage a tres basse temperature dans l'aluminium de baute purete', A cta Met 14 (1966) 1325 Panseri, C., Ceresara, S., Federighi, T. 'Recovery of aluminium cold-worked by compression at 78 K', Nuovo Cimento 29 (1963) 1223 Benoit, W., Bays, B., Grandchamp, P. A., Vittoz, B., Fantozzi, G., Perez, J., Gobin, P. 'Modulus effects and Hasiguti peaks on two coldworked fcc metals: gold and aluminium', J Phys Chern Solids 31 (1970) 1907 Gondi, P., Valdr'e, U., Grilli, A. 'Work hardening and recovery phenomena in aluminium', Nuovo Cimento 21 (1961) 834 Lytton, J. L., Westmacott, K. H., Potter, L.C. 'Recovery behavior of pure aluminium and selected body-centred cubic metals', ML TDR-64 189, WPAFB, Ohio, (1964) Masing, G., Raffelsieper, J. 'Mechanische Erholung von Aluminium - Einkrystallen', Z Metallk 41 (1950) 65 Perryman, E. C.W. 'Recrystallization characteristics of superpurity base A1-Mg alloys containing 0.5 pct Mg', TransAIME 203 (1955) 369 Lytton, J. L., Meyers, C. L., Tietz, T. E. 'The effect of "elastic and plastic strain on tensile flow stress recovery of aluminium', TransAIME 224 (1962) 1250 Tietz, T. E., Meyers, C. L., Lytton, J. L 'Recovery of tensile flow stress of aluminium and A l - l % Mg alloy' TransA1ME 224 (1962) 339

291