2. UJli

I. FA-TA-74037

FA

AD TA

74037

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4" THERMOMECHANICAL TREATMENTS ON HIGH STRENGTH

Al-Zn-Mg(-Cu) ALLOYS

LOAN COPY: RETURN TO AFWL TECHNICAL LIBRARY

KIRTLAND AFB, N. M.

December 1974

Approved for public release; distribution unlimited.

fOR" A# 20081120345 fir ^

Pitman-Dunn Laboratory

PENS*

-*i I.ibnur A I' 'A,.

1

U.S. ARMY ARMAMENT COMMAND FRANKFORDARSENALj

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*. TITLE (and Subtitle)

THERMOMECHANICAL TREATMENTS ON HIGH STRENGTH Al-Zn-Mg(-Cu) ALLOYS

5. TVPE OF REPORT ft PERIOD COVEREO

Technical research articl^ S. PERFORMING ORG. REPORT NUMBER

7. AUTHOR).)

E. Di Russo, M. Conserva, F. Gatto, * and H. Markus

8 CONTRACT OR GRANT NUMBERS)

9. PERFORMING ORGANIZATION NAME ANO ADORESS

Frankford Arsenal Attn: SARFA-PDM Philadelphia, PA 19137

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*E. Di Russo, M. Conserva, F. Gatto, Istituto Sperimentale dei Metalli Leggeri, Novara, Italy.

19. KEY WORDS (Continue on reverse mlde II neceeeary and Identity by block number)

Aluminum Alloys Thermomechanical Treatments Precipitation Hardening

Fracture Toughness Fatigue Stress Corrosion

20. ABSTRACT 'Continue on I eery and Identity by block number) An investigation was carried out to determine the metallurgical properties of Al-Zn-Mg and Al-Zn-Mg-Cu alloy products processed according to newly developed Final Thermomechanical Treatments (FTMT) of T-AHA type. The results show that these cycles can be utilized to produce wrought products of high purity Al-Zn-Mg(-Cu) alloys characterized by equivalent toughness and ductility and much higher strength than conventionally processed commercial purity

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20. Abstract - Cont.

materials. Based on transmission electron microscopy studies, it was found that such improved behavior of FTMT material is attributable to the superposition of hardening effects, from aging precipitation and from dislocations. Preliminary stress-corrosion and fatigue tests indicate that these properties are not substantially influenced by T-AHA thermo- mechanical process. Further work is needed in this area, in order to better understand the directions to follow for developing better alloys.

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Thermomechanical Treatments on

High Strength Al-Zn-Mg(-Cu) Alloys E. DI RUSSO, M. CONSERVA, F. GATTO, AND H. MARKUS

An investigation was carried out to determine the metallurgical properties of Al-Zn-Mg and Al-Zn-Mg-Cu alloy products processed according to newly developed Final Thermomechanical Treatments (FTMT) of T-AHA type. The results show that these cycles can be utilized to produce wrought products of high purity Al-Zn-Mg(-Cu) alloys characterized by equivalent toughness and ductility and much higher strength than conventionally processed commercial purity materials. Based on transmission electron microscopy studies, it was found that such improved behavior of FTMT material is attributable to the superposition of hardening effects, from aging precipitation and from dislocations. Preliminary stress-corrosion and fatigue tests indicate that these properties are not substantially influenced by T-AHA thermomechani- cal process. Further work is needed in this area, in order to better understand the directions to follow for developing better alloys.

JA.ECENTLY, increasing recognition has been given to thermomechanical treatments (TMT) as important tech- niques of improving the properties of metallic mate- rials. The use of such treatments has been limited primarily to iron base alloys and therefore it has been generally accepted that the term TMT refers to those treatments which involve deformation prior to or during phase transformations.1 In addition to this definition all of the metallurgical processes based on the combination of plastic deformation and heating, but not involving al- lotropic changes, have been reported as mechanical- thermal treatments (MTM).2 However, other types of alloy/treatments exist which cannot be classified under the strict definition of TMT but for which the term MTM is too general. Such a situation is represented by processes comprising plastic deformation and precipi- tation phenomena in aluminum alloys. The structural changes involved in these treatments cannot be con- sidered as simple rearrangements of dislocations be- cause the interaction between the lattice defects and the decomposition of the solid solution is characterized by complex synergistic effects. For this reason it seems reasonable to use the term TMT for those treat- ments on aluminum alloys that involve combination of plastic deformation and precipitation processes to give new strength-structure relations, even though aluminum alloys do not exhibit polymorphic transformations and as such do not strictly fall within the definition of TMT as reported in Ref. 1.

In addition, in aluminum alloys, different types of precipitation phenomena occur. The principal ones are: the high temperature precipitation of ancillary elements, which is usually produced during the thermal proces- sing of the cast material, and the low temperature pre- cipitation of solute rich particles, which occurs during the age-hardening stage. In order to give better clarity to the work, for thermomechanical cycles of alumi- num alloys involving precipitation of ancillary ele-

E. DI RUSSO and M. CONSERVA are Research Metallurgists and F. GATTO is Director, the Istituto Sperimentale dei Metalli Leggeri, Novara, Italy. H. MARKUS is Chief, Materials Engineering Division, Pitman Dunn Laboratory, Frankford Arsenal, Philadelphia, Pa. 19137.

Manuscript submitted June 26. 1972.

ments we have proposed the use of the term ITMT (Intermediate Thermomechanical Treatments), whereas when low temperature age-hardening phenomena are involved, the term FTMT (Final Thermomechanical Treatments) is suggested.

In the present work, the application of FTMT to high strength Al-Zn-Mg(-Cu) alloys has been investigated; our goal was to establish the conditions in which the dislocations introduced by plastic deformation interact most favorably with the age-hardening processes to produce materials that have improved properties com- pared to conventionally aged materials.

1) APPROACH

An overall view of the tensile properties of the most widely used commercial Al-Zn-Mg(-Cu) alloys is shown in Fig. 1. It can be seen from this figure that with the standard processed alloys, yield strength (YS) levels of 90,000 psi can be approached with elongation values of about 10 pet. Recently, considerable effort has been directed towards improving the strength of wrought Al-Zn-Mg(-Cu) alloys and some investigators produced these alloys with YS values that were higher than 100,000 psi, Fig. 1. Various combinations of strength and elongation have been achieved by Di Russo by con- trolling the solidification conditions, chemical compo- sition and the homogenization, solution treatment and aging cycles. Flemings and coworkers' obtained simi- lar results on alloys that were solidified under control- led conditions, completely homogenized and then cold worked (H) after normal solution (T) and aging (A) treatments. The use of different combinations of aging and work hardening was tested by Mercier et a/.5 using a process based on plastic deformation after complete T6 heat treatment, followed by a partial recovery treat- ment. Studies were also made by Pavlov et al. in which the solution treatment was followed by warm deformation (Hc) and then by artificial aging (A).

From Fig. 1 it can be seen that the increase in strength is always accompanied by a decrease in elong- ation which, at the highest YS level, drops to a very low value. This is more apparent for the treatments involving combinations of deformation and aging. Neglecting the T-AH treatments, in which no signifi-

MlDnrtl TDAUCA^TIflMC \: i i l II1JC

1000 PS \1 131

1 (4) 15)

\ •J 7001

7178

V 7075

{typical values U.S. specifications)

7039!.

" 0 5 10 15 20 E V. 25

Fig. 1—Yield strength values as a function of elongation shown by commercial and experimental alloys of Al-Zn-Mg(-Cu) system. The numbers in parentheses indicate the reference numbers.

cant relation exists between the two factors A and H, it is well known that the sequence T-HA in Al-Zn- Mg(-Cu) alloys, which involves plastic deformation after solution treatment, leads to a reduction in the response to subsequent artificial aging.7'8

In fact, the two parameters H and A are competitive so that properties of materials in the T-HA state are not substantially different from those of the TA state.5"11

However, developments in the knowledge of the aging mechanisms and of the microstructural features of the Al-Zn-Mg(-Cu) alloys,12"18 have suggested that it should be possible by means of more sophisticated TMT to modify plastic deformation and precipitation phenomena to get improvements in strength without drastic loss of ductility. To this end it was felt that the inhibiting effect that dislocations exert on the aging reactions could be substantially reduced by an initial aging step at a low temperature which would be given prior to the plastic deformation and the final aging step (T-AHA cycles).

This type of TMT has not been extensively studied but there are some data on Al-Cu and Al-Mg-Si alloys.8'19'20 Our preliminary investigations in 1967 on Al-Zn-Mg alloys21-23 showed that TMT involving solution, quenching, aging, working, aging (T-AHA), markedly increased the ultimate tensile strength (UTS) and YS values of these alloys compared to the conven- tional tempers, while still maintaining adequate ductil- ity.

More recently, extensive research at our laboratory on Al-Zn-Mg(-Cu) alloys confirmed that these new TMT (T-AHA)* produce alloys that have structures with im-

•It. Pat. No. 886,185. U.S. Filed 14.4.70 Scr. No. :8,5I4.

proved properties compared to conventionally heat treated materials.24"31 In addition T-AHA cycles were shown to have a beneficial effect on other heat treat-

able aluminum alloys.32'33

This paper presents a summary of the data obtained on the effect of TMT on the structure-property relation- ships in Al-Zn-Mg(-Cu) alloys. In more detail, se- lected results showing the tensile properties, notch and fracture toughness, fatigue and stress-corrosion behavior of different alloy/treatments are presented and discussed. Other papers will be published in the near future dealing with the more fundamental aspects of TMT.

2) EXPERIMENTAL PROCEDURE

a) Material Selection

The basic guidelines used for selecting the materials to study were: i) to have Al-Zn-Mg(-Cu) alloys with a wide range of solute contents. Thus, the alloys studied were 7005, 7039, and 7075. The composition of the al- loys investigated and the types of wrought products used are shown in Table I; ii) to have industrially pro- duced and laboratory produced alloys. The industrially produced alloys were all of commercial purity and were prepared as either slabs or ingots (alloy Nos. 1, 2, 4, 5). The laboratory produced alloys were prepared as both high purity (alloys Nos. 3, 7) and commercial purity (alloy No. 6) in the form of semicontinuous dc cast 110 mm diam ingots using controlled solidification conditions.31'35 The typical microstructures of the lab- oratory produced ingots in both the as-cast state and after homogenizing at 450°C for 8 h + 480'C for 15 h are shown in Fig. 2. The dendrite arm spacing (DAS) of the alloys range from 20 to 30 n. Due to the small DAS, the homogenization treatment used was sufficient to dissolve all of the eutectic network present in the high purity alloy ingots. In the alloy of commercial purity (alloy No. 6 : c. p. 7075) some of the secondary compounds containing the impurity elements, Fe and Si. combined with Al and Cu, or Al, Cu, and Mg remain undissolved. In the industrially produced commercial purity alloys, the solidification rates are lower and hence many coarse undissolved phases, containing iron and silicon.remain after the homogenization treatment. These microstructural differences play an important role in the behavior of materials. In fact, it has been shown that the presence of undissolved phases in 7000 series alloys has a detrimental effect on some im- portant properties;3'36"39 iii) to have materials with different grain structures. For this reason, sheet, plates, and extrusions were considered. Typical micro- structures of the wrought products in the fully heat treated condition are shown in Fig. 3. It can be seen that the sheets have a completely recrystallized struc- ture, the plates have a partially recrystallized struc- ture, i.e. subgrain structure, and the extrusions have a fine subgrain structure. It can also be seen that the high purity alloys have a much lower volume percent of undissolved second phase particles than do the com- mercial purity alloys.

b) Cycle Selection

The cycles were selected so that the effect of the deformation (H) in relation to the artificial aging stage (A) could be investigated.

The cycles used were T-A, T-AA, T-HA, T-HAA,

I I 1A lini HUE A APHII 107} MCTAI i Tier, i r A i TOAWCArTinwc

Alloy

7005 commercial purity 7039 commercial purity 7039 high purity

7075 commercial purity

7075 commercial purity 7075 commercial purity

7075 high purity

No. Zn Pet

Mg Pel

5.07 1.04

4.01 2.78

4.09 3.04

Cu Pel

Table I. Materials Investigated

Composition

Ci

Pel \ln

Pel

0.18

5.49 2.46 1.69 0.13 0.20

5.78 2.10 1.68 0.19 0.05

5.74 2.60 1.63 0.22

5.68 2.50 1.60 0.22

Pel Ti Pel

0.14 0.22 0.12

0.14 0.23 0.10

0.33 0.14 0.045

Fe Pel

0.12

0.14

0.012

Si Pel

0.09

0.09

0.017

0.074 0.28 0.20

0.057 0.12 0.056

0.005 0.22 0.13

0.010 0.0016 0.005

Wrought Products

4 mm-thick sheets and 70 by 7 mm extruded bars produced from an industrial dc cast ingot 4 mm-thick sheets produced from an industrial dc cast ingot 2 mm-thick sheets and 15 mm-diam extruded bars produced from laboratory dc cast 110 mm-diam ingot 2 mm and 5 mm-thick sheets and 15 mm-diam extruded bars produced from an industrial dc cast ingot 35 mm-thick industrial plale

2 mm-thick sheets, atrd 50 by 12 mm extruded bars produced from laboratory dc cast 110 mm- diam ingot 2 mm-thick sheets, 15 mm-diam and 50 by 12 mm extruded bars produced from laboratory dc cast 110 mm-diam ingot

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Fig. 2—Typical microstructures of the dc cast 110 mm-diam ingots of the 7039 h.p. (A), 7075 h.p. (B), and 7075 c.p. (C) alloys: (a) as cast: magnification 50 times; electrolytic etching in 2 pet HF; (ft) as cast: magnification 475 times; etching in 0.5 pet HF; (r) homogenized at 450°C for 8 h + 480°C for 15 h; magnification 475 times; etching in 0.5 pet HF.

METALLURGICAL TRANSACTIONS VOI IIMK4 APPii 1233 m;

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Fig. 3—Typical microslructures of some wrought products of 7039 and 7075 alloys, after full heat treatment (T6) observed on the longitudinal section. Magnification 475 times; etching: 70159 — in 55 pet H3PO4 at 50°C; 7075— in 25 pet HNC)3 at - 70°C. (a) 7039 c.p.. 4 mm-thick sheet; (ft) 7075 c.p.. 5 mm-thick sheet; (c) 7075 c.p., 35 mm-thick plate; (d) 7039 h.p., 2 mm-thick sheet; (e) 7075 h.p., 2 mm-thick sheet; (f) 7075 c.p., 16 mm-diam extrusion.

T-AH, T-AAH, T-AHA, and many combinations of A and H were selected. The values of H ranged from 5 to 50 pet reduction in thickness for sheets and from 5 to 20 pet reduction in thickness for plates and extrusions. The deformation was applied by rolling or drawing both at room temperature and in the range 130c to 220CC with preheating times ranging from 30 s up to 20 min. The first aging step, on the basis of hardness tests,40'41

was carried out at 100'C for 10 h for the ternary alloys and at 105rC for 6 h for the quaternary alloy.

c) Testing

All specimens were tested in tension in the longitudi- nal direction and for the 35 mm thick plate (alloy No. 5) and the 50 by 12 mm extrusions (alloy Nos. 6 and 7) also in the long transverse direction. At least two specimens were tested for each treatment.

The fracture toughness tests were carried out on 50 by 12 mm extruded laboratory produced 7075 (alloy Nos. 6 and 7) in the long transverse direction.

The specimen used to determine Kc had a single side notch and was similar to the specimen proposed by Sullivan.42'43 All of the values of Kc reported here are the average of 3 measurements.

The stress corrosion tests were performed on 35 mm thick plate of commercial purity 7075 (alloy No. 5) in the short transverse direction. "C" ring specimens were taken from the plate in the T-A, T-AA. and T-AHA tempers. The specimens were loaded to 50 and 75 pet of the yield strength and tested in alternate immersion in a 3.5 pet NaCl solution Each cycle consisted of 1 h (10 min immersion and 50 min exposure to air) and was carried out in a thermostatically controlled chamber at

20° ± 1°C. The test ended when cracks became visible. In the cases where no cracks were visible after 300 h, the samples were removed after various times up to a limit of 1000 h, unloaded and examined for the presence of plastic deformation.

If plastic deformation had occurred, optical micro- scopy was used to determine if stress corrosion cracks were present. Also, in doubtful cases and always after 1000 h testing, optical microscopy was used. At least 5 samples were tested for each temper in order to determine the specific time or the time interval where stress corrosion cracking had occurred.

The fatigue tests were carried out at room tempera- ture on cantilever beam specimens using a rotating beam testing machine operating at 11,500 rpm. Both smooth and notched specimens (Kt = 4,8) of 15 mm diam extruded bars of high purity 7039 and 7075 (alloy Nos, 3 and 7) were tested in the T-AA (ternary alloy), T-A (quaternary alloy), and T-AHA tempers. Also commercial purity 7075 (alloy No. 4) was tested in the T-A temper. The metallographic aspects of the mate- rials were observed by electron microscopy, to establish the relationships between properties and structure.

3) RESULTS AND DISCUSSION

a) Tensile Properties

The tensile properties of the most interesting treat- ments are given in Tables II to VI. It can be seen that with FTMT (T-AHA cycles) all the materials tested had significant increases in UTS and YS, with an acceptable decrease in elongation, compared to the values obtained with traditional cycles.

ill 11 M P A A PR II METALLURGICAL TRANSACTIONS

Table II. Longitudinal Tensile Properties Shown by 4 mm-Thick Sheets of 7005 and 7039 Alloys of Industrial Production in Several Tempers

Aging Treatment

1st Step

Deformation

2nd Step

Tensile Properties

Alloy Type Temperature Deg Pet UTS, psi YS, psi E.Pct

100°Cfor200h - - — - 57,000 49,500 16.5 100°Cfor 10 h - - - 130°Cfor22h 55,000 48,000 17.0 100°Cfor lOh rolling i t. 20 100°Cfor lOOh 63,000 59,500 12.4

7005 100°Cfor 10 h rolling r.t. 20 H5°Cfor 15 h 63,000 59,000 11.5 commercial purity 100°Cfor 10 h rolling it. 20 130°Cfor8h 62,000 57,000 1 1 1 No. 1 100°Cfor 10 h rolling t.t. Alt 100°Cfor50h 67,000 62,500 8.7

100°Cfor 10 h rolling r.t. 40 115°Cfor 10 h 65,500 61,000 8.8 100°Cfor 10 h rolling r.t. 40 130°C for4h 64.000 60,000 8.6 100°Cfor200h rolling r.t. 20 - 67.000 65,000 5.5 100°Cfor200h rolling r.t. 40 - 71,000 69,000 4.9

100°Cfor200h - - - - 65.000 55.500 16.0 100°Cfor 10 h - - 130°C for 22 h 63,500 55,000 15.4 100°Cfor 10 h rolling r.t. 20 100°Cfor60h 74,000 69,500 10.2

7039 commercial purity No. 2

100°Cfor 10 h rolling r.t. 20 115°C for 15 h 72,500 67,000 11.2 100°Cfor 10 h 100°Cfor 10 h

rolling rolling

i t.

r.t. 20 20

130°Cfor6h 100°Cfor30h

71,500 77,000

66,000 71.000

9.8 s.l

100°Cfor 10 h rolling r.t. 20 115°Cfor 10 h 76,000 70.000 7.8 100°Cfor 10 h rolling r.t. 20 130°Cfor4h 74,000 69,000 8.0 100°Cfor200h rolling r.t. 20 - 76,500 74,000 5.6 100°Cfor200h rolling r.t. 40 82,000 80,000 5.0

0 5 10 15 20 E V. 25

Fig. 4—Yield strength values vs elongation of sheets of 7005, 7039, and 7075 alloys in several thermal and thermomechani- cal tempers.

Sheet exhibits the best combination of strength and ductility, with increases in UTS and YS from 15 to 30 pet and decreases in elongation of 2 to 8 points, equiva- lent to 10 to 50 pet compared to the traditional tempers T-A and to processes of basic T-AH type. In Fig. 4 the shaded areas represent the range of values ob- tained with T-AHA cycles and the lower curve is de- fined by the values relative to the T-A, T-AA, T-AH, and T-AAH tempers for each alloy.

Less of an increase in strength was found in plate and especially in extrusions, which, because of their

micrograins and oriented structure, have high tensile properties even in the traditional tempers. It should be pointed out that other treatments comprising strain and aging do not give favorable strength-ductility combinations.

For example, the T-AH or T-AAH cycles, at a level of strength equivalent to that obtained by the T -AHA cycle, show a loss of elongation from the T6 temper of 10 to 12 points which is equivalent to a reduction of 70 to 80 pet.

If the deformation is applied before artificial aging (T-HA cycles) the peak strength values are lower, other conditions being equal, than those obtained with the T-AHA cycle, Table V. Also, if H is not very high, the final peak values are lower than those obtained with conventional thermal treatments. The intermediate deformation increases the aging rate during the final artificial aging step; for example, in 4 mm thick sheet of 7005, after an initial aging at 100CC to 10 h followed by 20 pet deformation, the peak strength during subse- quent aging at 100°C is reached in about 100 h, whereas in the normal isothermal aging treatment the peak strength does not occur before 200 h, Table II. It has been found that the deformation should not exceed 30 pet for sheets, 20 pet for plates, and 10 pet for extrusions since higher degrees of deformation cause only a slight increase in strength accompanied, however, by a large decrease in elongation.

Regarding warm deformation, it can be applied with no deleterious effect in the range 130' to 170°C for the ternary alloys (preheating times 10 to 20 min) and in the range 170° to 220°C for the quaternary alloys (pre- heating times 2.5 to 15 min up to 190CC and 30 s to 2.5 min for higher temperatures). As can be seen for 7039 alloy, Table III, for equal final properties, the strength- ening achieved by the last aging step is higher in the case of warm deformation than for the cycle with H applied at room temperature. For 7075, the use of warm deformation, in addition to producing increased plasticity of the alloy, results in a better combination

METALLURGICAL TRANSACTIONS VOLUME 4. APRIL 1973 1 137

Table III. Longitudinal Tensile Properties of 4 mm-Thick Sheets of 7039 C.P. Alloy (No. 2) in Several Thermomechanical Tempers, with Deformation Given at

Various Temperatures

Aging Treatment

Tensile Properties Deformation

2nd Step 1st Step Type Temperature Deg Pet UTS, psi YS. psi E, Pel

100°Cfor 10 h - - - 130°Cfor22h 63.500 55,000 15.4 100°Cfor 10 h - - 54.000 30,000 22.9 100°Cfor 10 h rolling r.t. 20 - 66,000 58,000 9.0 100°Cfor 10 h rolling r.t. 20 115°C for 15 h 72,000 66,500 10.2 100°Cfor 10 h rolling 130°C :n 64,500 55,000 10.0 100°Cfor 10 h rolling 130°C 20 115°C for 15 h 72,000 66,500 10.0 100°Cfo. 10 h rolling 150°C 20 62,500 55.000 10.5 100°Cfor 10 h rolling I50°C 20 115°C lor 15 h 71,500 65,500 9.7

100°Cfor 10 h rolling 160°C 20 59,500 52,000 9.5 100°Cfor 10 h rolling 160°C 20 115°C for 15 h 70,000 65,000 10.1 100°Cfor 10 h rolling 170°C 20 61,000 52,000 9.7 100°C for 10 h rolling 170°C 20 115°C for 1 5 h 69.500 62,500 9.4

Table IV. Tensile Properties of Sheets (Longitudinal Direction) and Plates (Long Transverse Direction) of 7075 C.P.

Alloys of Industrial Production in Several Tempers

Material

Aging Treatment

Tensile Properties

1st Step

Deformation

2nd Step Alloy Wrought Product Type Temperature Deg Pet UTS. psi YS, psi E, Pet

120°C for 24 h - - - 83.000 73,500 13.1 105°Cfor6h rolling r.t. 20 105°Cfor50h 95,500 91,500 7.8 105°C for 6 h rolling r.t. 20 120°Cfor8h 92,000 88,000 8.0

5 mm-thick sheet 105°Cfor6h rolling r.t. 50 120°Cfor4h 98,000 94,000 4.8 120°Cfor24h rolling r.t. 20 95,500 92,000 4.0 !05°Cfor6h rolling 160°C 20 120°Cfor8h 92,000 86,000 7.7 105°Cfor6h rolling 190°C 20 120°Cfor8h 91,000 86,500 8.2

7075, No. 4 120°C for 24 h - - 81,000 71.000 16.7 105°Cfor6h rolling r.t. 10 120°Cfor 13 h 89.000 83,000 13.0 105°Cfor6h rolling r.t. 20 120°Ct'or 7h 92,500 88,000 10. 1

2 mm-thick sheet l05°Cfor6h rolling 170°C 10 120°Cfor 13 h 88,500 86,000 12.3 105°Cfor6h rolling 170°C 20 120°C for 6 h 93,000 88,000 10.5 105°Cfor6h rolling 190°C 10 120°C for 13 h 90,000 85,000 11.7 105°Cfor6h rolling I90°C 20 120°Cfor5h 92.000 89.500 10.7

120°C for 24 h - - 78,500 71,000 8.0 7075, No. 5 35 mm-thick plate 105°Cfor6h rolling 175°C 10 120°C for 13 h 83,000 78,500 5.0

105°Cfor6h rolling 180°C 20 120°C for 7 h 83.000 79,500 5.0

of tensile properties than is obtained by the use of room temperature deformation, Fig. 5 and Table IV.

In all of the T-AHA cycles studied the high purity 7039 and 7075 sheets have good toughness even though the notch toughness slightly decreases with increasing degrees of deformation. On the other hand, in com- merical purity 7075, greater than 10 pet deformation caused a significant decrease in notch toughness, Table V. The effect of FTMT on the strength of both com- mercial purity and high purity alloys is the same. It has been found that high purity 7075 sheet has higher ductility, and especially, notch toughness, than its commercial purity counterpart, independent of the type of treatment used, Table V. In practical terms, this means that by using high purity alloys, materials pro- duced in T-AHA temper at 15 pet higher yield strength have equal or higher elongation and reduction in area values, higher notch strength and notch toughness as compared to the corresponding commercial purity al- loys in the T-A temper. It should also be pointed out, Tables V and VI, that there are some differences in the

tensile properties of the industrially produced and laboratory produced commercial purity alloys.

Generally, the alloys prepared in the laboratory have better combinations of strength and ductility for a given temper; this is probably due to their smaller dendrite arm spacings.

b) Fracture Toughness

The results of fracture toughness tests along with the corresponding tensile properties of samples of laboratory produced 7075 in the T-A and T-AHA tempers are shown in Table VII. It can be seen that although FTMT generally causes a decrease in tough- ness, h.p. 7075 in the T-AHA state (H = 10 pet) does have comparable elongation values, 17 pet higher K values and 12 pet higher yield strength values than commercial purity 7075 in the T-A temper.

c) Stress Corrosion Resistance

The short-transverse stress corrosion resistance and the corresponding long transverse tensile prop-

1 138 VOLUME 4, APRIL 1973 METALLURGICAL TRANSACTIONS

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RI190°Cx15 rrwi)*He B*A120°C I Wlw

R(220°C x 15 rwn> • He 10 * A120°C x 6 hr

fH»'A120°C x 13hf

1RI2208C x 30 MC ) * He 10 * A120 "C x 2*f»

-Hk; JL 0 20 170 160 190 200 210 220 *C

fa)

_fc

R|170 'C x 15 mm) * He 20 • A120 8C x E hr

hf R|190°C«2.5m.nl*Hc20*At20oC* 6hr

R(220aC x 25 min) • He 20 • A120 °C x 10 hr

fRI170*Ci1S mml*Hc20'A12O°Cx 6ht

I-Att5*Cx6hfWRIT90oCx15 mm)*Hc20'A120BC x 5hr

1R(220'CX15 mm]*Hc20*A120«C 1 6hi

|H20*A120*Cx7hf

lRt22O»Cx30 Me)*He20*A12O°Cx Bhr T-A10S°C x6hr

To

(b) Fig. 5—Tensile properties of 2 mm-thick sheet of 7075 c.p. alloy (No. 4) subjected to thermomechanical treatments of T-AHA type, with H = 10 pet (a) and 20 pet (b), given at dif- ferent temperatures. The times of final aging at 120°C correspond to peak strength.

erties of the 35 mm thick plates of commercial purity 7075 in various T-AHA cycles, as well as the T6(T-A), T73(T-AA), and a stress corrosion resistant T-AA temper,44 are presented in Table VIII.

It can be seen from this table that T-AHA cycles with the second aging step at low temperatures (speci- mens d and e in Table VIII) have higher strength than the T6 material but do not show any improvement in stress corrosion resistance.

In the T-AHA cycle where the second aging step is carried out at a higher temperature (specimen / in Table VIII), at strength levels equal to or slightly higher than those of the T6 temper, there is a marked improvement in the stress corrosion resistance over the T6 temper. However, it does appear that the com- plex dislocation-precipitation structure obtained with FTMT has only a small effect on the stress corrosion resistance since in the T-AA and T-AHA states which have the same aging treatments (specimens c and / in Table VIII), the stress corrosion resistance is the same.

d) Fatigue

The tensile properties of the 7039 and 7075 tested are shown in Table VI. The S-N curves for the fatigue tests on the 15 mm diam extruded bars of high purity 7039 and commercial and high purity 7075 are shown in Figs. 6 and 7, respectively. The S-N curve of smooth specimens of 7039 T-AA does not show a well defined limit while that of the T-AHA material has a higher slope up to about 107 cycles and has a well defined asymptote. The fatigue strength at 5 x 108 cycles for the materials in the T-AA and T-AHA tempers are 30,000 psi and 28,500 psi, respectively.

In the notched specimens of 7039, the T-AA and T-AHA tempers show fatigue strengths which at 5 x 108

cycles are 14,000 psi and 9,000 psi, respectively. As in the case of the smooth samples, material in the T-AHA temper has a higher slope at short times and a more pronounced limit which becomes stable at a lower number of cycles. The effective notch coefficients (K eff) which is the ratio between the limit (or the fatigue strength at 108 cycles) for the smooth specimens and that for the notched ones are: T-AA temper K eff = 2.10 and T-AHA K eff = 3.33. Thus, it appears that T-AHA makes the alloy more notch sensitive than the T-AA temper.

In smooth specimens of commercial purity 7075 there is no definite limit observed up to 109 cycles. The high purity 7075 in the T-A temper has a well defined limit after about 107 cycles and has a fatigue strength at 5 x 108 cycles of 35,000 psi. This value is higher than the fatigue strength of the commercial purity material at 5 x 10s cycles. In the T-AHA material the S-N curve is regular up to about 5 x 107 cycles and with increased cycles, there is a sharp decrease in the fatigue strength with a discontinuity phenomenon, or "knee" effect45'46

which is probably associated with differences in the

hp7039 Ki-4.8

X v-

^; •*H$

I0J 10' 10' 10" 107 cycles 10s 10J 10' 105 106 107 cycles 10'

T-A100°Cxl0hr .AI30°Cx22hr T- A100°Cx 10 hr .He 12 7.. A115°Cx 30hr

Fig. 6—S-N curves for 15 mm-diam extruded rods of 7030 h.p. alloy (No. 3) with various treatments, obtained on smooth (kt = 1) and notched (kt = 4.8) specimens.

METALLURGICAL TRANSACTIONS VOLUME 4, APRIL 1973- 1 139

Table V. Longitudinal Tensile Properties {Smooth and Notched (kf = 10.2) Specimens) of 2 mm-Thick Sheets of 7039 and 7075

Laboratory Produced Alloys, in Several Tempers

Aging Treatment

Tensile Properties

1st Step

Deformation

2nd Step Alloy Type Temperature Deg Pet US, psi "1 S, psi E, Pet RA, Pet MS, Psi NTS, YS

7075 high purity No. 7

120°C for 24 h 84,500 72,000 17.8 26.9 86,500 1.20 l05°Cfor 6h rolling 175°C 10 120°Cfor 13 h 90,000 84,000 16.2 27.2 94,500 1.12 I05°C for6h rolling 175°C 20 120°Cfor6h 94,000 89,000 13.0 19.4 92,400 1 in

rolling r.t. in 1 20°C for 24 h 82,000 69,000 16.0

7075 120°C for 24 h 87,000 75,000 15.3 28.0 81,000 1.08 commercial purity IOS°Cfor6h rolling 175°C 10 l20°Cfor 13 h 93,000 86,500 10.4 22.9 79,000 0.91

No. 6 IOS°Cfor6h rolling 175°C 20 120°Cfor6h 95,500 91,000 10.0 16 8 75,000 0.82

7039 high purily No. 3

100°Cfor 10 h 130°Cfor22h 66,000 58,000 11 8 38.5 70,500 1.21 100°Cfor 10 h 100°('for 10 h

rolling rolling

I55°C I55"C

It) 20

H5°Cfor28h 115°Cfor 15 h

72,000 72,500

67,000 68,500

12.4 10.6

32.4 30.9

76.000 78.000

1.14 1.14

100°CI'or 10 h rolling 15S°C 40 115°Cfor9h 75,500 72,000 10.5 18.6 79,000 1.10

Table VI. Tensile Properties of 15 mm-Diam Extruded Rods of 7039 and 7075 Alloys of High and Commercial Purity in Several Tempers

Aging Treatment Tensile Properties

1st Step Type

Deformation

2nd Step UTS, psi YS, psi Alloy Temperature Deg Pet E, Pet

7039 com me re I.I] IOO°Cfor 10 h - - 130°Cfor22h 82,000 77,000 8.2 purity, No. 2 l00°Cfor 10 h drawing r.t. 20 11 5°C for 1 5 h 87.000 85,500 4.8

7039 high 100°Cfor 10 h - - 130°Cfor22h 83,500 77,000 7.5 purity, No. 3 100°Cfor 10 h drawing 155°C 12 115°C for 30 h 85,000 82,000 6.7

7075 commerical 1 20°C for 24 h - - - 101,000 94,000 5.7 purity. No. 4 10S°Cfor6h drawing 180°C 10 120°C for 10 h 106,000 105,000 3.4

7075 1 20°C for 24 h - — 120°Cfor24h 100,500 96,500 6.2 high purity l05°Cfor6h drawing 175°C 5.5 120°Cfor24h 104,000 101,500 5.6 No. 7 105°C for 6 h drawing 175°C 12 120°Cfor 13 h 105,000 103,000 4.3

Table VII. Long-Transverse Tensile and Fracture Toughness Properties of 50 by 12 mm Extruded Bar of 7075 Laboratory Produced Alloys, in Several Tempers

' Aging Treatment

Tensili ! Properties

1st Step

Deformation

2nd Step

Fracture Toughness

Alloy Type Temperature Deg Pet UTS, psi YS. pis 1 , Pet Kc

Index, psi Vi"

7075 commercial purity No. 6

7075 high purity No. 7

120°Cfor24h 105°C for 6 h 105°Cfor6h

120°Cfor24h 105°C for 6 h 105°Cfor6h

rolling rolling

rolling rolling

175°C 175°C

175°C 175°C

10 20

10 20

120°Cfor 13 h 120°Cfor6h

120°C for 13 h 120°Cfor6h

87.(100 93,000 92,500

89,000 93,500 95,000

78,500 86,000 89,000

81,500 87,500 89,500

8.8 5.6 5.0

[0.6 8.1 7.8

40,500 28,000 27,000

57,000 47,000 37,500

mechanism of either crack initiation or crack propa- gation at different stress levels. For this reason it is difficult to define a fatigue limit and thus, one can only describe a length of time at a certain load level. In this sense, the alloy does not fail up to about 5 x 108

cycles under a stress of about 25,000 psi. The notched specimens of 7075 in the T-A temper

show a typical curve with the "knee" effect and a well defined fatigue limit at 107 cycles of about 11,000 psi. This limit appears to be independent of the degree of purity of the alloy. The experimental values for the S-N curve of notched specimens of 7075 in the T-AHA temper follow two curves which appear to be sufficiently different to suggest that this behavior is not due to

usual scatter in the data. Up to 108 cycles, the upper curve is quite different from the T-A material, while the lower curve is typical for notched specimens with a well defined limit at 106 cycles. The effective notch coefficients for 7075 in the T-A temper and the T-AHA temper are 3.22 and 1.77, respectively. These values show that there is a marked reduction of the notch ef- fect in the samples produced with FTMT cycles, show- ing quite a different behavior of 7075 from 7039.

The results on the whole suggest that T-AHA cycles produce a decrease in the fatigue resistance of the ternary alloys, in particular in notched specimens. In the case of the quaternary alloy, T-AHA cycles lead to more reduced fatigue properties on smooth specimens,

1140 VOLUME 4, APRIL 1973 METALLURGICAL TRANSACTIONS

PSI

SO

40

30

20

10

n

•\ h.p 7075

\ .

- c p. 7075 ""^^^^

' ""Mill ' " • "LIUI

'

V h p. 7075

\ *•

to

"max 1000 PSi h.p.7075

K| =4.8

* 1 i. c.p.

i X. 075

_

v. \ - H ̂

-ll -

h.p.7075 K|»4.8

\ V \ s • . ^.

- ^-L_iS **•

cycles

h.P.

Fig. 8—Transmission electron micrograph of 2 mm-thick sheet of 7075 h.p. alloy (No. 7) in T-H10 pet A120°C for 24 h temper. Magnification 19,200 times.

101 10' K)5 106 107 cycles 105 10s 10' 10b 10s

T-AI20°Cx24hr T-A105°C x 5 hr.Hc 12'/. •

Fig. 7—S-N curves for 15 mm-diam extruded rods of 7075 alloy (No. 7) with various treatments, obtained on smooth (kt = 1) and notched (kt = 4.8) specimens. For comparison S/N curves for similar samples of 7075 c.p. alloy (No. 4- dotted line) in T-A120°C for 24 h are reported.

while causing an improvement in notched samples. However, the discrepancies in the data do not permit unambiguous interpretation especially since, in respect to other metallurgical properties, the relationship be- tween the fatigue processes and microstructure is not well known. Thus it appears a difficult task to explain the S-N curves in terms of dislocations and precipita- tion distributions.

e) Transmission Electron Microscopy

Transmission electron microscopy examinations show that, similarly to what has been reported elsewhere for the ternary alloys,23 the improvement in strength pro- duced by FTMT with respect both to conventional ther- mal treatments and to other processes based on differ- ent combinations of H and A, can be attributed to a better hardening type structure in the FTMT samples.

Specifically, Figs. 8 to 10 show at low magnification that high purity 7075 sheets have uniform dislocation structure in the T-H10 pet A120°C for 24 h and T-A 105°C for 6 h H10 pet A120°C for 13 h states in con- trast to the defect free matrix in the T-A sample. At higher magnification, the T-HA sample, aged to peak hardness, has a coarse precipitation of r\' (or if) par- ticles with respect to a fine distribution in both the T-A and T-AHA cycles, Figs. 11 to 13. In other words, it appears that T-AHA samples have the optimum combination of dislocations and precipitate structure for effective strengthening.

A typical feature of the T-AHA samples is the pres- ence of wide slip bands overlapping the uniform dis-

Fig. 9—Transmission electron micrograph of 2 mm-thick sheet of 7075 h.p. alloy (No. 7) in T-A105°C for 6 h H10 pet A120°C for 13 h temper. Magnification 48,000 times.

Fig. 10—Transmission electron micrograph of 2 mm-thick sheet of 7075 h.p. alloy (No. 7) in T-A120°C for 24 h temper. Magnification 47,500 times.

METALLURGICAL TRANSACTIONS VOLUME 4. APRIL 1^7.3 1141

Fig. 11—Transmission electron micrograph of 2 mm-thick sheet of 7075 h.p. alloy (No. 7) in T-H10 pet A120°C for 24 h temper. Magnification 190,000 times.

Fig. 12—Transmission electron micrograph of 2 mm-thick sheet of 7075 h.p. alloy (No. 7) in T-A105°C for 6 h H10 pet A120-C for 13 h temper. Magnification 192,000 times.

1

lift ••SL,

' * ' ' * .'

location structure and characterized by dense tangles of dislocations hardly resolvable also at high magnifi- cations, Fig. 14(a),(ft). The tendency to form such slip bands has been found to be greater in samples which have a high degree of intermediate deformation and in wrought products, such as extrusions, which are con- stituted by a subgrain structure. The presence of these coarse slip bands may play a role in the loss of ductil- ity in FTMT material.

4) CONCLUSIONS

The increase in strength without drastic loss in ductility achieved in medium and high strength Al-Zn-Mg(-Cu) alloys by the application of T-AHA type FTMT is quite significant. This increase in strength is probably due to a superposition of the strengthening effects from dislocations and from pre- cipitation. The most beneficial effect of FTMT occurs in high purity recrystallized wrought products, i.e. sheet; in these materials, the YS and UTS can be in- creased by 20 pet with the same or better ductility and notch and fracture toughness compared to commercial

"*"

f *

Fig. 13—Transmission electron micrograph of 2 mm-thick sheet of 7075 h.p. alloy (No. 7) in T-A120°C for 24 h temper. Magnification 192,000 times.

Fig. 14—Typical aspect of coarse slip bands observed on 4 mm-thick sheets of 7005 c.p. alloy (No. 1). treated according to a T-AHA cycle. Magnification (a) 16.800 times; (b) 48.000 times.

^^i- METALLURGICAL TRANSACTIONS

Table VII Tensile (Long Transverse) and Stress-Corrosion (Short Transverse) Properties Determined on 35 mm-Thick Plate of 7075 C.P.

Alloy (No. 5) of Industrial Production, in Several Tempers

Aging Treatment Stress Corrosion Properties

Type 1st Step

Deformation

Type Temp. Deg Pet 2nd Step

Tensile Properties

UTS, psi YS, psi l.l'cl

a) 120°Cfor24h

Applied Load

Pet YS

75

78,500 71,000 8.0

50

bl 105 C tor 6h 170 C for 8 h 75,000 67,500

105°Cfor6h 150°Cfor8h 79,000 72,000 8.0

105°Cfor6h rolling 175°C 10 120 C for 13 h 83,000 78,500 5.0

105 C for 6 h rolling 180 C 20 120°Cfor7h 83,000 79,500 5.0

105 C for 6 h rolling 180 C 10 150 C for 8 h 79,000 73,500 7.9

75

50

75

50

75

50

75

50

75

50

Time of Test, h

48 to 72

160 to 300

-1000

350 to 500

800 to 1000

40 to 80

100 to 250

-24

40 to 90

400 to 500

800 to 1000

Optical Microscopy Observation

Intergranular cracks (some* times passing through the thickness) which start from the exposed tensioned surface

Dispersed shallow pits (--0.15 mm); intergranular cracks (not passing through) starting from the bottom of some pits or directly from the exposed tensioned surface

Dispersed pits rather deep ( 0.30 mm); intergranular cracks (not passing through) starting from the bottom of some pits. Some transgranular cracks in the corroded zones

Dispersed pits more or less deep ( 0.25 mm); intergranular cracks (not passing through) starting from the bottom of some pits or directly from the exposed tensioned surface

Dispersed pits rather deep (-0.35 mm); intergranular cracks (not passing through) starting from the bottom of some pits

Intergranular cracks (some- times passing through) starting from the exposed tensioned surface

Some shallow dispersed pits (-0.1 mm); intergranular cracks (sometime passing through) starting in general directly from the exposed tensioned surface

Intergranular cracks (some- times passing through) starting from the exposed tensioned surface

Dispersed pits more or less deep (-0.25 mm); intergranular cracks (not passing through) starting from the bottom of some pits or directly from the exposed tensioned surface

Dispersed pits rather deep (-0.35 mm); intergranular cracks (not passing through) starting from the bottom of some pits

conventionally aged material. FTMT does not produce significant variation in the typical stress corrosion resistance but may cause a slight decrease in the fatigue resistance; however, this work does show that FTMT in combination with the use of high purity alloys does give a method of producing materials with much higher strength and the same or better secondary prop- erties than conventionally available materials.

It should be pointed out that there are no problems in the application of FTMT cycles to simple shaped wrought products, especially sheet and plate. The work-

ing operations can be carried out at relatively high temperatures, so that the application of the deformation itself is facilitated. Also, there are no close tolerance limitations on the deformation temperatures or on the aging treatments, so that FTMT cycles seem to be quite compatible with good production control standards.

ACKNOWLEDGMENTS

The authors acknowledge the valuable assistance provided by Mr. M. Buratti in performing this work.

METALLURGICAL TRANSACTIONS .^..^•a

They also indebted to the following ISML technicians: ing. Belvedere (plastic transformation); dr. Bracale (preparation of ingots); dr. Farina (stress-corrosion tests); Mr. D. Morri (tensile and fracture toughness tests); ing. Screm (fatigue tests).

This work was sponsored by the Italian Ministry of Defense and the U.S. Department of the Army, under a cooperative research and development project.

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vol. 20, pp. 770-74. 7. B. Chalmers: Prog, in Mater. Set, vol. 10, ed., p. 151, Pergamon Press, Oxford.

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10. J. I. Polmear and P. Scott-Young: /. Inst. Metals, 1958. vol. 87, pp. 65-70. 11. G. M. K. Sarma and T. R. Anantharaman: Trans. Indian Inst. Metals, 1965,

vol. 18, pp. 151-54. 12. E. Di Russo: Final Rep. on contract number DA-91-591-EUC 3782, ISML,

Novara, July 1966. 13. E. Di Russo: AUuminio, Nuova Met., 1967, vol. 36, pp. 348-63. 14. E. Di Russo: Final Report No. AD-676I07 on contract number DAJA37-67-0-

0387, ISML, Novara, April 1968. 15. W. Gruhl: Aluminium. 1962, vol. 38, pp. 775-79. 16. W. Gruhl: Z. Metallic., 1964, vol. 60, pp. 577-82. 17. D. O. Sprowls: U.S. Patent No. 3,198,676, August 1965. 18. P. L. Mehr, E. H. Spuhler, and L. W. Mayer: Alcoa Alloy 7075T73, Alcoa

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30. M. Buratti: Report No. 70/20 616, ISML, Novara, September 1970. 31. I.S.M.L.: Final Report on CTSD-32 Contract, April 1971. 32. E. Di Russo and M. Conserva: Italian Patent No. 886 185. 1971, U. S. Patent

Case No. 28514, 1970. 33. M. Buratti, M. Conserva, and E. Di Russo: Report No. 71/21 904 ISML, Novara,

November 1971. 34. G. Bracale and A. Alti: Report No. 69/19 563, ISML, Novara, July 1969. 35. G. Bracale: Report No. 71/21 170. ISML. Novara. March 1971. 36. S.N.Singh and M.C. Fleming: Trans. TMS-AIME, 1969, vol. 245. pp. 1811-19.

37. H. W. Antes, S. Lipson, and H. Rosenthal: Trans. TMS-AIME, 1967, vol. 239, pp. 1634-42.

38. H. W. Antes and H. Markus: Met. ling. Quart., 1970, vol. 10, pp. 9-11. 39. J. H. Mulherin and H. Rosenthal: Met. Trans., 1971, vol. 2, pp. 427-32. 40. M. Conserva, E. Di Russo, M. Buratti, and F. Gatto: Mater. Sci. Eng., 1973,

vol. 11, pp. 103-12. 41. M. Buratti and E. Di Russo: AUuminio, Nuova Mel., 1970, vol. 39. pp. 69-75. 42. A. H. Sullivan: Mater. Res. Stand., 1964, vol. 4, pp. 20-24. 43. D. Morri: Report No. 70/20 565, ISML, Novara, 1970. 44. C. Arcolin, M. Buratti. and E. Di Russo: Report No. 71/21 547. ISML, Novara,

July 1971. 45. F. Gatto: Rev. Met., 1958, vol. 60, pp. 1085-90. 46. D. Morri: AUuminio, 1958, vol. 27, pp. 387-90.

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Administration Attn: Code RRM Federal Building #10 Washington, DC 20546

Director National Bureau of Standards Attn: Tech. Info. Div. Washington, DC 20025

Director U.S. Atomic Energy Commission Attn: Document Library Germantown, MD 21403

1 Attn: AFML-Technical Library 15

Defense Documentation Center (12) Cameron Station Alexandria, VA 22314

Metals and Ceramic Information Center

Battelle Memorial Institute 505 King Avenue Columbus, OH 43201

Mr. Robert H. Brown 1411 Pacific Avenue Natrona Heights, PA 15065

Dr. Robert S. Busk 3606 Windsor Court Midland, MI 48640

Mr. J. B. Hess Kaiser Aluminum & Chemical Corp. Aluminum Division Research Center for Technology

P.O. Box 870 Pleasanton, CA 94566

Prof. G. R. Irvin University of Maryland College Park, MD

Dr. Schrade Radtke International Lead Zinc Institute 292 Madison Ave. New York, NY 10017

Prof. M C. Flemings Department of Metallurgy Materials Science Mass. Institute of Technology Cambridge, MA 02139

Mr. Carson L. Brooks Reynolds Metal Company 4th & Canal Streets Richmond, VA 23219

Mr. Harold Hunsicker Aluminum Company of America Alcoa Technical Center Alcoa Center, PA 15069

Prof. J. Kenig Drexel University Phila, PA 19104

Dr. E. J. Ripling Materials Research Laboratory, Inc 1 Science Road Glenwood, IL 60425

Prof H. Rogers Drexel University Phila., PA 19104

Frankford Arsenal:

1 Attn: Commander's Reading File AOA-M/107-B

1 Attn: Technical Director/107-1

1 Attn: J. Mc Caughey, CE/107-1

1 Attn: Director, PD/64-4

1 Attn: G. White, PD/64-4

1 Attn: PDM-P/513

1 Attn: PDM-E/64-4

1 Attn: PDM-A/64-2

1 Attn: Director, MD/220-1

1 Attn: MDM/220-1

1 Attn: MDA/220-3

1 Attn: MDS/220-2

1 Attn: FI/107-B

1 Attn: MT/211-2

1 Attn: MTT/211-2

10 Attn: PDM/64-4

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Frankford Arsenal - Cont'd

2 Attn: TSP-T/51-2

2 Attn: TSP-L/51-2

1 - Reference copy 1 - Circulation copy

1 Attn: PA/107-2

1 Attn: GC/28-1

1 Attn: QAA-R/119-2

Printing & Reproduction Division FRANKFORD ARSENAL Date Printed: 19 December J.974

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