NRL Report 6607

Metallurgical Characteristicsof High Strength Structural Materials

Twelfth Progress Report

R. J. GOOnD, R. W. HusBK, D. G. HowE, R. W. JUDY, JR.. T. W. CROOKER,

E. A. LANGE, C. N. FRZED, AND P. P. PUZAK

Strength of Metals BranchMetallurgy Division

September 1967

7 NAVAL RESEARCH LABORATORYWmhington, D.C.

CONTENTS

Abstract ..................... .................... vProblem Status ................ .................... vAuthorization ................. .................... v

INTRODUCTION ................ ..................... 1

HIGH STRENGTH TITANIUM ALLOYS ......... ........... 3

Fracture Toughness Index Diagram forTitanium ................. .................. 4

Fracture Toughness Studies ....... ........... 4

Heat Treatment Study on Ti-6A1-2Cb-lTa-0.8Mo 6

Feasibility Study on Composite TitaniumAlloy Plate .......... .................. ... 13

THE EFFECTS OF HEAT TREATING ENVIRONMENTAL CONDITIONSON THE STRESS-CORROSION-CRACKING RESISTANCE OFSEVERAL TITANIUM ALLOYS ....... .............. .. 16

Materials and Procedures ..... ............ .. 16

Experimental Results ....... .............. .. 18

STRESS CORROSION CRACKING OF SOME TITANIUM ALLOYS . . 28

Current Results of SCC Study .... .......... .. 29

Fractographic Study of Fatigue Fractures . . . . 57

LOW CYCLE FATIGUE CRACK PROPAGATION IN A201B, A302B,AND A517F PRESSURE VESSEL STEELS ................ 59

Experimental ........... .................. .. 63

Test Specimen and Materials .... .......... .. 63

Apparatus and Procedure ..... ............ .. 63

Results and Discussion ..... ............. ... 68

Fatigue Crack Propagation in Air ......... ... 68

Fatigue Crack Propagation in AqueousEnvironments ......... ................ .. 71

PAGE BLANK I

CONTENTS (Contd)

Comparisons with Full-Size Vessel TestResults . . . . . . . . . . . . . . . . . 77

Summary and Conclusions . . . . . . . . . . . 80

CORRELATION OF TWO FRACTURE TOUGHNESS TESTS FORHIGH-STRENGTH STEELS . . . . . . . . . . . . . 81

Materials and Procedure .*. . . . . . . .... 82

Analysis of Data ....... . .. . . 82

ALUMINUM ALLOYS. ....... .. ....... 88

Stress Corrosion Cracking . . . . . . . . . . 88

Fracture Toughness Index Diagram forAluminum, ... & a a * *. . . . . . ...* ae 95

HIGH STRENGTH STEELS . . . . . . . . . . . . . . . 97

Fracture Toughness Index Diagram for Steels . 97

Drop-Weight Tear Tests (DWTT) of 5Ni-Cr-Mo-VSteel Plate [HY-130(T)] . . . . . . . . . . 99

Full Thickness DWTT of 2-3-in. Thick Steels .103

Drop Weight '½ar Test Study of 2%Mn-2%Ni WeldMetals Developed for Fabricating HY-130(T)Steel Weldments . .. .. * e * a * * * * * .105

Charpy V (Cv) Fracture Toughness of HY-130(T)Plates and 2%Mn-2%Ni Weld Metals . . . . .110

DWTT of "New" 9Ni-4Co-0.20C Alloy High StrengthSteels . . . . . . . . . . . . . . . . . . ..117

DWTT of 9Ni-4Co and Maraging Stoel WeldMetals. . . . . . . . . .. . .. . .. . 121

REFERENCES . . . . . . . . . . . . . . . . . o 126

iv

ABSTRACT

A progress report covering the research studies in highstrength structural metals conducted during the periodJuly 1966 through January 1967 is presented. The reportincludes fracture toughness studies on some new titaniumalloys as well as results of an initial feasibility studyaimed at raising the optimum strength and toughness limitsfor titanium alloys through an approach involving thicksection composites. The results of salt water stress-corrosion-cracking studies on a number of titanium alloysare described including a study concerning the effect ofvacuum heat treatment on stress-corrosion-cracking resis-tance of titanium alloys. Fatigue crack propagationstudies in air and in salt water on the pressure vesselsteels A302B, A201B, and A517F are discussed; and theresults compared to the actual service performance ofthese same materials in PVRC program studies. Prelimi-nary fracture toughness correlation diagrams are presentedfor high strength steels based upon fracture mechanics andengineering test methods. The latest versions of theFracture Toughness Index Diagrams for steels, titaniumalloys and aluminum alloys, a usual feature of this reportseries and which are based upon engineering test methods,are presented and the significance of the featuresbriefly discussed. Fracture toughness studies on thickplates of 5Ni-Cr-Mo-V, and a "new" 9Ni-4Co steel arereported. In addition, the results are presented forsimilar studies on welds of 2Mn-2Ni in 5Ni-Cr-Mo-Vsteels and welds of 9Ni-4Co, 12Ni-3Cr-3Mo and 17Ni-2Co-3Mocompositions.

PROBLEM STATUS

This is a progress report; work is continuing.

AUTHORIZATION

NRL Problem F01-17; Project S-4607-11894NRL Problem M01-05; Projects RR-007-01-46-5405,

and SF-020-01-01-0724NRL Problem M01-18; Projects RR-007-01-46-5420,

SF-020-01-05-0731, andENG-NAV-67-1

NRL Problem M03-O; Projects RR-007-01-46-5414,and SF-020-01-01-0850

NRL Problem M04-08B; ARPA 878

Manuscript submitted June 26, 1967.

METALLURGICAL CHARACTERISTICS OFHIGH STRENGTH STRUCTURAL MATERIALS

LTwelfth Progress Report]

INTRODUCTION

This report is the twelfth in the series of status reportscovering the U.S. Naval Research Laboratory MetalurgyDivision's long-range Advanced High Strength StructuralMetals Program. This program is concerned with determiningthe performance characteristics of high strength metals andis directed to developing the necessary information to pro-vide "guideline" principles for the metallurgical optimiza-tion of alloys, for processing and fabrication techniquesand for reliable failure-safe utilization of these materialsin large complex structures. Fracture toughness aspects ofthe failure-safe design problem are being studied withrecently developed and conventional, long-establishedengineering test methods. The subcritical crack growthcharacteristics of the high strength metals are studiedunder conditions of low cycle fatigue. Environmentaleffects on subcritical crack growth and the applicationof fracture mechanics techniques for the spectrum of materialsunder investigation are also being studied.

The drop-weight tear and explosion tear tests and the Charpy Vnotch (for steels in particular) have been used to determinethe fracture toughness characteristics of steels, titaniumalloys, and aluminum alloys over wide ranges of yieldstrengths and have made possible the development of frac-ture toughness index diagrams for these materinas. Thesediagrams which have been featured in this Progress ReportSeries, provide guideline information for purposes ofdesign, alloy development, specification, and quality con-trol. The latest versions of these diagrams for steels,titanium alloys and aluminum alloys are presented in thisreport.

Preliminary fracture toughness characteristics have beendetermined for the new alloys Ti-6A1-4V 2Mo, Ti-6AI-6V-2Sn-2Mo, Ti-4A1-3Mo-lV. Ti-2A1-4Mo-4Zr, and a new heat ofTi-6A1-4V. The results indicate that the Ti-4Al-3Mo-lValloy has considerable potential for high toughness in the120 ksi yield strength range although welding may be aproblem. A Ti-6A1-2Cb-lTa-O.8Mo alloy from a heat-treatment

study indicates a high level of fracture toughness below110 ksi yield strength provided the interstitial contentis kept low and proper heat-treatments are uRod.

Initial feasibility studies on the use of a compositeapproach to make high strength and higher toughnesstitanium plate material am presented. Explosive bondingand roll bonding techniques were employed to produce theplate. The results of the fracture toughness, metallo-graphic, and ultrasonic studies have provided considerableencouragement at this early phase of what is envisionedas a long-range effor, to provide a major advancement inthis high and ultrahigh strength metals technology.

The effects of vacuun heat treatment on imparting essentialimmunity to salt water stress-corrosion-cracking to other-wise stress-corrosion-cracking sensitive titanium alloysare ,.escribed. At the same time, the interstitial hydrogenlevel is reduceu to very low levels due to the vacuum heattreatment whereas similar heat treatments in inert envir-onments neither impart a similar high resistance to saltwater nor reduce the interstitial hycrogen to the samelow levels. The details are described for the Ti-8A1-lMo-lV,Ti-7AI-lMo-IV, Ti-6A1-4V, and Ti-7A1-2.5Mo alloys. Inaddition, the stress-corrosion-cracking characteristics ofa number of additional titanium alloy plate and weldmentsin salt water have been determined and the results arepresented along with the results of a fractographic studyof fatigue crack surfaces developed during preparation offracture mechanics type fracture toughness specimens. Insome cases, cleavage and quasi-cleavage was noted to bepresent, but it was not possible to conclusively associateit with the fatiguing operation or the initial crackextension phase in the fracture mechanics test.

Low cycle fatigue crack propagition studies in the A201B,A302B, and A517F pressure vessel steels in air and in saltwater are *.escribed in one section of this report. Sincethese steels represent pedigreed material obtained fromfailed pressure vessels, comparisons are made between thelaboratory test results and actual structural experiencewith these materials. It %as found that in air fatiguecrack grcwth rates in A201B and A302B steels follo- acommon third power relationship with total strain range;a fourth power relationship %as found for A517F. A2OIBpossessed the greatest tolerance for cyclic plastic strain.

2

All of the steels showed some sensitivity to salt waterbut no gross effects were observed. Comparisons betweenthe laboratory results and structural experience withthese steels were in agreement for the A201B and A302Bsteels.

The results of a preliminary study on steexs of therelationship between plane strain fracture toughness andthe results obtained with engineering type fracture tough-ness tests, in this case, the drop-weight tear test, arereported. A number of high strength steels were investi-gated using the single-edge notched, side-grooved tensionspecimen and a remarkably good correlation was establishedbetween drop-weight tear test energy for fracture and theplane strain fracture toughness KIc, the strain energyrelease rate bic and $Ic (plastic zone size indicator).These correlatioi,; peruiit one to make reasonable estimatesof Kic,& ic, and $Ic from nreasured drop-weight tear energyvalues.

The salt water stress-corrosion-cracking resistance ofthe aluminum alloys 7075-T7351, 7079-T6 and 7005-T63 isreported. The first twxo alloys appeared to be relativelysensitive to the environment in the "short transverse"direction whereas the 7005-T63 did not.

Fracture toughness studies of steel plates aad weldscharacterized by strengths ranging from about 130 to 200ksi yield strength have been conducted. The steels in-cluded 3/4 to 4-inch thick plates of 5Ni-Cr-Mo-V and1-inch thick plates of a "new" version of the tNi-4Co-0.XXCquenched and tempered steel. Producer fabricated manualand metal-arc inert gas welds of the nomijnal 2Mn-2Ni alloycomposition developed for welding the 5Ni-CL-Mo-V steel,and NRL fabricated, fully-automatic, tungsten-arc inertgas welds of the 9Ni-4Co-O.XXC alloy and l2Ni-3Cr-3Mo and17Ni-2Co-3Mo maraging steel weld metal compositions werealso examined.

HIGH STRENGTH TITANIUM ALLOYS(R. J. Goode and R. W. Huber)

Preliminary fracture toughness characteristics have beendetermined for the recently received 'e% alloys Ti-6A1-4V-2Mo,Ti-6A1-6V-2Sn-2Mo, Ti-4A1-3Mo-iV. Ti-2A1-4Mo-4Zr, and a newheat of Ti-rAA-4V in the as-received condition, additional

3

data are also pre.anted for several of the alloys after asolution annealing treatment in vacuum. In support of aheat treatment study being conducted by Reactive MetalsCorporation, full thickness fracture toughness propertieswere de-ermined for a Ti-6A1-2Cb-lTa-0.8Mo alloy afterheat treatment.

A new study concerning the properties and characteristicsof thick section composites of titanium alloys has beeninitiated and fracture toughness data has been obtainedfor several composites representative of initial attemptsat defining significant parameters and their control forcladding thick sections.

FRACTURE TOUGHNESS INDEX DIAGRAM FOR TITANIUM

The latest fracture toughness index diagram for titaniumis shown in Fig. 1 along with the drop-weight tear test(DWTT) data points for all materials studied. The featuresof the diagram and their significance have been describedin detail ii. earlier reports of this series and in Refer-ences 1-3. Briefly, below 1500 ft-lb DWTT energy materialswill propagate fractures through elastic stress regions;increasing levels of plastic strain are required for propa-gation of fractures for increasing DWTT energy values above1500 Ft-lb. The impingement of the elastic-to-plastictransition bond (1500-1700 ft-lb DWTT energy) with theoptimum materials trend line (OMTL) indicates that forthose materials having yield strengths above 140 ksi elasticstresses can be expected to sustain fracture propagation.Between 120-140 ksi yield strength (YS) only carefullyoptimized alloys will be capable of developing plasticstrains before fracture occurs. The region below120 ksi YS is characterized by a number of alloys thatrequire high levels of plastic strain for propagation offractures.

FRACTURE TOUGHNESS STUDIES

One-inch thick plates were recently received of the follow-ing alloys:

1. Ti-6A1-4V-2Mo (R-2) - a modification of the Ti-6A1-4Valloy

2. Ti-6A1-6V-2Sn-2Mo (R-4) - a modification of theTi-6A1-6V-2Sn alloy

3. A high purity Ti-4A1-3Mo-lV alloy (R-3)

4

-SB1-iA) Ag83N3 A AdtVH:) -O 30NVdI

0D 0

Fr

4 0

VA

O's

-0

<8

;"! . * ,* -

o o 0

to- o.........

) It 0 0

($87-1-I) A983N3 8V31 IH')13M dO8OFig. I The FTID for 1-inch thick titanium alloys

5

4. Ti-2A1-4Mo-4Zr (R-5) - a new Titanium Metals Cor-poration of America alloy

5. Ti-6A1-2Mo (R-l) - represents an additional heatof a promising high strength, high toughness alloy previ-ously studied.

The DWTT fracture toughness of the as-received materialwas determined for these alloys. In addition, a single DWTTspecimen from. the plates R-I,R-2, and R-4 were vacuum solu-tion annealed for 1 or 2 hours in a cold wall vacuum furnaceat temperatures judged to be just below the A transus of eachalloy (0 transus temperatures to be determined). The DWTTresults are shown in Table 1 along with the tensile dataprovided by the producer as well as those presently deter-mined at NRL. The DWTT energy data is plotted against theFTID features in Fig. 2 for the as-received material andsince the tensile data is available for the Ti-6A1-6V-2Sn-2Moalloy after the vacuum heat treatment. The NRL tensile datais used in this figure where possible. (It should be notedthat these data do not nece3karily represent the capabilitie2of these alloys but are related only to the beginning of anoptimization study on these materials.)

The combination of 2325 ft-lb DWTT energy and about 115ksi YS for the Ti-4A1-3Mo-lV alloy in the as-received condi-tion indicates that this alloy may have considerable potentialfor high toughness in the 120 ksi YS range. However, thehigh content of beta forming elements may lead to difficul-ties in fusion welding. An indication that such a problemmay exist was observed when an electron beam, bead-on-platetype of weld decreased the DWTT energy to 573 ft-lbs. Moreextensive, MIG type welding studies on this alloy as wellas on the others is scheduled to begin as soon as likecomposition welding wire on order is received.

HEAT TREATMENT STUDY ON Ti-6Al-2Cb-lTa--0.8Mo

The Eleventh Quarterly Report (7), contained data con-cerning the fracture toughness of DWTT specimens of aTi-6Al-2Cb-lTa-0.8So alloy (1-inch plate, HT No. 292555).The specimens represented a variety of heat-treated condi-tions and were evaluated in the DWTT at NRL for ReactiveMetals Corporation in support of their Navy sponsored heat-treatment study on this alloy. Afterwards the broken halveswere returned to Reactive Metals Corporation to providematerial for determination of the related tensile andCharpy V (Cv) properties. Thesedata have now been obtained

6

TABLE 1

Tensile and Fraoture Toagmeu Properties of Sam Tita•,imIlloys (Plates R- twocud •& ..&) . _ _ -

2% Y.S. UTS hag R. A.(kol) (kni)(%) (M)

Allot DIM _ P!RM EM PROD EM PMQ NDL HF) ERL

R-1 1W 1662 As ree'd 13U.8 114.2 126.9 126.1 13.5 9.8 29.3 20.3

6Al-M WR 1601 As ree'd 111.9 117.5 124.9 126.6 14.5 10.5 29.8 27.1

WR 2146 V.A. 1 Hr18500 1PHe Cooled

R-2 WR 931 An reo'd 130.9 125.3 11.5.6 141.1 13.0 10.0 24.4 19.0

6A1-4V-2Mo RU 1052 As reo'd 126.1 126.0 1U3.4 141.0 12.5 10.0 21.6 21.0

RH 811 V.A. 1 Hr 185001PHe Cooled

R-3 RW 2325 As rec'd 114.6 - 122.2 - 16.5 - 40.6 -

tA1-3Mo-lV WR 2325 An reo'd 115.5 - 123.5 - 15.0 - 38.8 -

R-4 R3 514 As ree'd 139.4 134.1 165.0 161.0 7.0 6.5 8.5 10.0

6AI-6V-2Sn-294o WR 514 As ree'd 141.6 138.2 168.6 161.3 6.0 6.5 10.0 10.9

WR 681 V.A. 1660 0 7 2 Hr - 126.8 - 148.1 - 10.0 - 10.0He Cooled

R-5 WR 2325 As reo'd 104.3 - 115.t - 20.0 - 51.2 -

2A1-O4o-hZr MW 2498 As rec'd 94.2 - 111.8 - 20.5 - 43.8 -

47L

TITANIUM ALLOYSONE INCH THICK PLATES

CODES50 0 0 ST RO N G W EA K u_

- Ti-6AI-2 Mo"I� g I Ti-6AI-4V-2Mo

_ 0 ® @ Ti-4A,-3Mo-IV>_ 4000 - I Ti-6AI-6V-25n-2 Mo 60-90

OM0 @ Ti-2A1-4Mo4Zr

'= 3000 - 40-60

1:2000o 15-40o

0a-o 0 5-20

010 I I i lI ,

80 90 100 110 120 130 140 150 160 1000 PSII I

40 50 60 70 TONS /INI - . I * I I,

60 70 80 90 100 110 KG/MM2

YIELD STRENGTH

Fig. Z - Fracture toughness characteristics of titanium alloy platesR-1 through R-5 in the as-received condition (plus vacuum 1660 Fheat treatment of R-4)

I

81

and are shown in Table 2 along with the chemical analysis ofthe alloy and the DWTT determinations. In additioi, similardata from another heat of Ti-6A1-2Cb-ITa-0.8Mo from the samestudy, which contains slightly less Cb and twice as michoxygen (0.122 versus 0.058%), is contained in this table.

The DWTT energy data, plotted against the FTID features,are shown in Fig. 3 for both heats of material. The numbersby the data points correspond to the heat treatment numbersin Table 2. Several general observations can be made fromthis data:

1. In general, this alloy is characterized as havinggood fracture toughness in 1-inch thicknesses below 110 ksi.yield strength.

2. High oxygen promotes lower fracture toughness(point 9 compared to point 6 - the 50OF increase in P anneal-ing temperature wouldn't account for the large drop in DWTTenergy).

3. The higher oxygen promotes greater strengthening(points 9 compared to 6, compared to 7; and 11 compared to8 - again the 50OF 0 annealing temperature differencewouldn't be expected to account for the significant strengthdifferences).

4. Aging at 1100 0 F for 2 hours increases the yieldstrength over that of the annealed material but at the sametime causes a significant drop in the DWTT energy. Thisindicates a potential welding HAZ problem may exist for thisalloy and should be investigated in the DWTT for shorteraging times and different temperatures.

5. Rapid cooling after annealing increases the strengthof the material substantially; its effect on fracture tough-ness is not clear cut from the data, but it appears to bebeneficial.

6. Charpy V energy at -80°F is not a very sensitiveindicator of changes in fracture toughness in the range ofstrength and toughness of concern here.

Additional DWTT samples of another heat of Ti-6AI-2Cb-lTa-0.8Mo from the same study were tested. The DWTT data 4,are shown in Table 3 for the different heat treated conditions.

9

TABLE 2

HEAT TREATMENT OF Ti-6A1-2Cb-lTa-0.8Mo 1-INCH PLATE

Al Cb Ta Mo Fe C N 0 Beta% % % % % % % % Transus

Heat 292555 6.2 2.40 1.00 0.74 .06 .02 .006 .058 18450 F+50

R.M.I. NRL

TENSILE AND IMPACT PROPERTIES

Drop-WeightHeat Test UTS' 0.2%YS' Elong' RA1 Cv Impact' Tear Energy

Treatment Direction (ksi) (ksl) (%) (%) -80*(ft-lb) +32*F(ft-lbb)

As Rolled L 119.6 101.8 10.8 23.1 23.5 2266T 123.0 105.0 12.5 33.6 23.0 2146

1600°F-lHr-AC L 118.2 99.2 12.5 28.0 29.2 2443T 122.4 107.4 12.5 31.9 32.2 2846

18150F-lPr-AC L 118.2 97.7 12.5 31.3 32.5 2560T 122.2 104.0 13.0 33.0 28.2 2266

1815 0 F-lHr-WQ L 126.0 101.2 13.0 29.0 27.8 2560T 129.0 108.4 11.8 30.7 32.2 2443

1815*F-lHr-WQ L 127.0 108.9 10.5 20.8 25.5 20861100*F-2Hr-AC T 135.5 118.9 10.5 22.1 29.0 1905

1900*F--1Hr-AC L 123.2 101.4 11.8 24.6 33.0 2618T 123.2 103.4 12.2 27.7 28.2 1966

1900°F-lHr-WQ L 131.1 110.2 12.8 31.0 30.5 2265T 128.1 108.0 8.8 23.8 26.8 2266

1900*F-lHr-WQ L 135.3 117.1 9.5 24.8 24.0 1723'.100 0F-2Hr-AC T 135.4 116.2 9.5 21.2 23.2 2146

Al Cb Ta Mo Fe C N 01 1 It It It T It

Heat X-3454 6.0 1.98 0.92 0.80 .05 .03 .006 .122

1950°F-lHr-AC L 127.6 105.5 12.2 27.6 24.8 1784T 128.5 106.3 12.5 26.4 25.5 --

1950*F-lHr-WQ L 135.2 115.0 11.5 27.3 22.5 --

T 136.6 116.6 10.8 19.4 24.0 --

1950 0 F-iHr-WQ L 137.1 119.8 12.5 25.4 22.0 11731100F-2Hr-AC T 139.0 121.1 9.2 16.0 23.2 --

'Average of two specimens

10

(SG1-LA) S3nflvAk!DN3N3 A Ad8HD J~O 3!DNVU q fl

0 0

S U)

(D 1 0

w z

OD 00 -

0A

In r') (D

(S8-IA 0S33?V1 HIMd-

Fig.3 - racure oughesscharcterstis ofTi-~Al-Cb-Tq.M

alloys~~ frMR1ha-tetnWtd

11 ¾

0.

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Cd C*4 CI t- t-

FU r 4-4 H

C4.

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w r44 r- UIr-4 I I I

r-4 Pi a

0 0

0

Until the tensile data and chemical analysis is obtainednot much can be said concerning this material except thatthe 2000OF solution anneal appears to be above the optimumannealing temperature for obtaining high fracture toughness.

FEASIBILITY STUDY ON CO(MPOSITE TITANIUM ALLOY PLATF

A long range study has been initiated on the crackpropagation resistance and mechanical property chiracter-istics of thick section composite titanium alloy plate.The scheme is to clad a low strength, high fracture tough-ness alloy on to both surfaces of a high strength, lowfracture toughness core material. Some of tne beneficialeffects of the cladding that might be expected are:

1. Promotion of extensive plastic flow at the surfacesof the composite plate which would provide a "dragging anchor"effect on a crack moving through the core material; hopefully,this anchor effect would tend to arrest what might otherwisebe a self-propagating crack in the core material.

2. Protection of Lne core material from stress-corrosion cracking in aqueous and other agressive environ-ments.

3. Provide high strength plate material which ishighly resistant to fatigue crack propagation from surfaceflaws.

The ultimate aim ol this study is to provide a signifi-cant advancement in the useful strength and fracture toughnesslimits for titanium alloys in complex structural applications,i.e., move the OM1TL in Fig. I to higher levels of strengthand toughness.

Through the cooperation of E. I. DuPont Company testpieces of a ELI grade (0.12 oxygen minimum) T'-6AI-4V alloywere "Deta Clad" (explusion bonded) o,- bolth sides %ith a1 16-inch and a 1/4--inch thick Ti-3.5AI lloy -,j form a platecomposite 1-inch thick. In addition, .i 2-inch thick plateof the Ti-641-4V alloy was "eta Ciad" with 1,/2-inch thickTi-3.5AI plate on ei;ch surface and then thc cý•.n-posite plate,as hot rolled t.,a final thickness of 1-inch. The finalcladding thicknoss for this plate %;.s about 3/16-inch

13

Three other titanium alloy compacts (20-inches by 20-inches) were built up by welding at the edges 1/2-inch thickcover plates of the Ti-3.5AI alloy to both faces of a 2-inchthick Ti-5AI-2V-2Mo-2Sn alloy, and 1/4-inch thick coverplates of Ti-3.5A1 to the faces of 2-inch thick Ti-5A1-2V-2Mo-2Sn and Ti-6AI-3V-lMo plates. These compacts were thenhot rolled by T.M.C.A. at 1700OF to a final thickness of1-inch and air-cooled off the rolling mill.

The bonding efficiency for the deta cladding processappeared very encouraging as determined by ultrasonic inspec-tion. Lack of bonding at one corner of the 3" explosionbonded compact was indicated. After hot rolling of thiscompact the whole composite was again ultrasonicallyexamined with this same area still showing a lack of bonding-the bonding on this plate was about 85% efficient. All ofthe T.M.C.A. roll bonded composites showed complete bonding.

DWTT specimens were cut from all these composite platesand tested; metallographic specimens were removed from sec-tions of the broken halves for further exami'iation of thebonded regions. The DWTT results for the ccmposite platesare shown in Table 4. The data indicates a decrease intoughness for the "Deta Clad" plate compared to the 811ft-lb DWTT obtained for the 1-inch ELI Ti-6AI-4V platematerial(5). This lowering of fracture toughness can beattributed to the cold working effect of the explosionshock loading in the cladding process. Vacuum annealingat 1700°F indicates that significant increases can beexpected in the fracture toughness of "Deta Clad" materialwhen the effects of cold work are removed. The DWTT ener-gies of all the hot rolled composites are significantlyhigher than the corresponding values obtained for thecore material alloy as 1-inch circular rolled plates(5)•Although all of these composites represent initial attemptsat defining the optimum cladding procedures and individualalloy characteristics in the bonding operations, theresults here appear encouraging with respect to the aimsof the study.

The tested DWTT specimens of the composites exhibitedevidence of some delamination at the fracture surfaces; inparticular with the "Deta Clad" plates. This delaminationappears to be associated with plastic deformation in theregion of the fracture: whether it occurred ahead of orbehind the advancing crack front is not known. Preliminary

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TABLE 4

FRACTURE TOUGHNESS OF 1-INCH THICK COMPOSITETITANILM PLATE

As-received DWii energy after VacuumAlloy DWTT energy Anneal at 1700OF

(ft-lb) (ft-lb)

Deta Clad Material(DuPont)

I. ELI Ti-6Al-4V (7-95)core material

2. I116-inch cladding 455 1418

3. 1/4-inch cladding 514 1966

,4. 1/2-inch deta cladon 2-inch core thenhot roll to 1-inch plate;3/16" resultant claddingthickness 2325 1844

Roll-bonded (T.M.C.A.)

1. Ti-5AI-2V-2Mo-2Sn (T-90)core material

a. I 16-inch cladding RW 2846WR 2206

b. 3 16-inch cladding RW 2184WR 2325

2. Ti-6A1-3V-IMo (T-93)

A. I 16-Inch ,'ladding RW 2676WR 1905

15

metallographic examination of the bond region of the hotrolled material shows a coherent interlace with little, ifany, grain growth occurring across it. The bond interfaceof the "Deta Clad" material in general shows a wave likeinterface; in many regions the interface shows mechanicalinterlocking between the two alloys as shown in Fig. 4.This interlocking is confined to within 2 or 3 grains ofthe bonded interface.

Large clad plafps are being processed to providematerial for considerably more exhaustive study in theinitial phase of this study.

*THE EFFECTS OF HEAT TREATING ENVIRONMENTAL CONDITIONSON THE STRESS-CORROSION-CRACKING RESISTANCE

OF SEVERAL TITANIUM ALLOYS"(D.G. Howe & R.J. Goode)

The effect of heat treating environmental conditionson the stress-corrosion-cracking (SCC) resistance of the.ery low interstitial (0.08 max oxygen and - 60-70 ppmhydrogen) alloys Ti-8A1-lMo-lV, Ti-7A1-lMo-lV, Ti-6A1-4V,and Ti-7A1-2.5Mo are being studied and the results of partsof the investigation have been reported in earlier reports(4-9).

MATERIALS AND PROCEDURES

Specimens, simi]ar in configuration to those used byBrown and Beachem (10) were obtained from approximately 1-inch thick plate material. All specimens were machined,sid3-notched and fatigued until the fatigue crack had propa-gated approximately 1/8 inch from the root of the notch.The specimens were degreased with Methyl Ethyl Ketone, washedwith acetone, and air dryed before heat treatment.

Heat treatments were conducted in either an Inconelmuffle in an electric furnace or a "cold wall" vacuum furnace.The cantilever SCC test procedure (10) was used to establishthe effect of heat treatment on the SCC resistance of thetitanium alloys in a 3.5% NaCl solution. In this procedurethe specimen is tested as a cantilever beam, with the pre-cracked zone surrounded by a plastic container holding the3.5% NaCi solution. A moment "M" is imposea upon the pre-cracked notch by weights at the end of the lever arm whichis clamped to the specimen, and the time is noted for final

7Thi-sresearch was supported by the~ Advanced Research Projects Agency

as part of the ARPA Coupling Program on Stress-Corrosion Cracking.

16

Ti-3.5AIClad

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Core

Fig. 4 - Bond interface between Ti-3.5AI "Deta"clad to ELI grade Ti-6A1-4V (T-95) showing ev-idence of mechanical locking. Mag. lOOX

17 -

fracture. If the specimen does not fail in the corrosiveenvironment at the preselected load, the moment is increaseduntil it does.

The torque "M" and specimen dimensions ars- then convertedto the stress intensity parameter KT through the use of equa-tions due to J. Kies (11). The minimum value of KI which sobserved to cause SCC is designated KIscc.

EXPERIMENTAL RESULTS

The data obtained in this investigation indicate thatfor the alloys studied the resistance to SCC can be signi-ficantly improved if a fine grain size is maintained throughprocessing and proper heat treatments are used. It is evi-dent that heat treatment in vacuum, followed by heliumcooling, is very beneficial as the alloys were relativelyinsensitive to attack by stress-corrosion-cracking underthe test conditions used.

The effects of vacuum and inert atmosphere heat treat-ments on the SCC resistance of the alloy Ti-7Al-lMo-.lV(T-88)were studied. The results of this study are shown in Tables5 and 6.

The solution anneals listed in Table 5 show the heattreatment given to the specimens before dividing them intothree groups for aging at 1200*F for 2 hours, under variousenvironmental conditions. The specimens were either agedin vacuum and helium cooled, aged in argon and air cooled,or aged in argon and water quenched. The environmentalconditions associated with the aging treatments at 1200OFhad little effect on the SCC resistance obtained as a resultof the solution annealing treatment. The material solutionannealed in vacuum and helium cooled showed high resistanceto SCC where the material heat treated at the same tempera-tures and time in an inert atmosphere (argon) followed byair cooling were susceptible to SCC under identical testconditions. It should be noted that the K, and timesassociated with the no-break situation in Tables 5 and 6indicate the next to last step loading condition. In someinstances, if the holding times were extended considerably,failure of the specimen might have occurred at the slightlylower K, values. The last column in Table 5 shows varia-tion of hydrogen content with different heat treatments.The Kiscc values shown in Table 5 were plotted against the

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interstitial hydrogen content present in the material afterheat treatment. This observed relationship is shown in Fig.5. It can be seen that the specimens which were susceptibleto SCC had a larger hydrogen contenit present in the structure.

The effects of different environmental conditionsassociated with 1 hour - 1700°F soluticn anneals are shownin Table 6. Four environmental conditions during and aftersolution annealing treatments we-e studied. All specimenslisted in Table 6 as beinE aged were heat treated togetherin vacuum at 1200°F for 2 hours and helium cooled. Thecooling rates as a result of these environmental conditionswere also determined and are illustrated in Figs. 6 and 7.

The SCC resistance of the Ti-7A1-iMo-lV alloy (T-88)resulting from different environmental conditions during aI hour - 1700*F solution anneal is illustrated in Fig. 8.The specimens solution annealed in vacuum followed by heliumcooling were highly resistant to attack by SCC. The speci-mens that were solution annealed in an argon or heliunatmosphere at the same temperature and time exhibited somesusceptibility to attack by SCC.

Tests conducted two months after the vacuum solutionannealing of SCC specimens taken from Ti-6AI-4V(T-27) andTi-7AI-2.5Mo(T-94) material showed both alloys to be highlyresistant to SCC. Recent tests to determine whether underlaboratory atmospheric environment the beneficial conditiondeveloped through vacuum solution annealing was a permanenteffect have shown that nine months after the heat treat-ments the Ti-6A1-4V and Ti-7A1-2.5Uo alloys are stillessentially immune to attack by stress-corrosion-cracking.The results of these tests are given in Table 7.

A preliminary summary of the Kiscc values obtained asa result of various heat treatments compared to the residualinterstitial hydrogen contents for all four of the alloysinvestigated is shown in Fig. 9. It is interesting to notethat for these alloys where the hydrogen contents were atvery low values the material was non-susceptible to stress-corrosion-cracking. Where hydrogen contents were greaterit can be seen that the alloys displayed varying degreesof susceptiblity to SCC. This would appear to indicatethat the interstitial hydrogen present in the alloy playsa role in the stress-corrosion-cracking reaction mechanism.

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S.C.C. RESISIANCE OF Ti-A-11Mo-IV RESULTING FROM DIFFERENTENVIRONMENTAL CONDITIONS DURING A I HOUR-1700F SOLUTION ANNEAL

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The data obtained to date in this investigation,indicate that a significant reduction of interstitialhydrogen content promotes high resistance to SCC. Indi-cated also is a possible threshold level of hydrogencontent below which the alloys are highly resistant toSCC and above which increased sensitivity to SCC develops.The hydrogen level of this threshold may be alloy dependentand may well fall at a level below the solubility limit ofthe alloy.

A perma, ent resistance to salt SCC would be dependentupon not having additional hydrogen absorbed into the struc-ture during actual service conditions,

*STRESS CORROSION CRACKING OF SOME TITANIUM ALLOYS(R. W. Judy, Jr. and R. J. Goode)

Stress-corrosion-cracking (SCC) studies have beenconducted on a wide variety of commercial and experimentaltitanium alloy plates. The tests wEre conducted by thecantilever method introduced by B.F. Brown (12).

The SCC resistance of a metal can be determined onlywhen certain conditions are met simultaneously; these con-ditions are a sufficient stress level and flaw size togetherwith an active environment. For a number of titanium alloys,an aqueous (salt water, distilled water, etc.) environmentis sufficiently active. The stress level and flaw sizenecessary to cause SCC to occur are very conveniently com-bined and expressed in terms of the stress intensity factorKI. A threshold level of stress intensity above which SCCwill definitely occur has been shown to exist for a numberof titanium alloys (13). This stress intensity level isdenoted by Kisc. Comparison of K icwith KI the stressintensity requ red for fracture in a r (designated "dry"),gives an indication of the relative resistance of an alloyto en'ironmental cracking. The formula used to calculateK values was the cantilever equation:

4.12M / & 3

KI BDý2

*This research was supported by the Advanced Research ProjectsAgency as part of the ARPA Coupling Program on Stress-CorrosionCracking.

28

where M = moment at test sectionB = specimen widthD = specimen thickness

S= 1 - a/Da = depth of flaw (notch and fatigue crack).

The specimens tested were 1/2 x 1 x 7 in. long witha machined notch and fatigue crack flaw located at itscenter in the WT or LT fracture direction (14). Thespecimens had shallow side grooves to insure plane strainconditions. A 3.5-wt-%-salt-water solution was used asthe aqueous environment.

CURRENT RESULTS OF SCC STUDY

The SCC resistance and mechanical properties (whereavailable) of several titapium alloys as one-inch-thickplate are shown in Table 8 and Figs. 11 through 27.Except as noted all the plates were in as-received (mill-annealed) condition and were low interstitial content(below 0.08% 02). Included in this set of data arecharacterization curves for several alloys in the as-received condition, a variety of heat treatments for aTi-7Al-lMo-lV alloy, and base plate data for materialtaken from heau-treated MIG weldments (weldment datashown in Table 9).

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Kisc 2

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where a is the critical crack depth (1).

Samples of Ti-7Al-lMo-lV taken from DWTT specimcns inthe as-received condition and solution heat treated at 1750OF

29

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and at 1850*F for one hour and then cooled in helium, weretested in the cantilever test; Kiscc:KIx values for thesewere 80:111, 92:122, and 90:117, respectively (Figs.These values indicated neithe•- an appreciable sensitivityto SCC nor an cbhvious effect of heat treatment in this case.However, specimens of base plate material cut from a MiGweldment of the sane material heat treated at 1850*F/l hr/He cooled indicated a pronounced sensitivity to SCC,Klsrc:K x being 42:118 (Fig. 15). The obvious differencesin gee tests were the difference in specimen width (1/2-inch wide for the specimens from DWTT specimens, 3/4-inchwide for specimens from the MIG weldment) and the differ-ence in the section si•'e being heat treated (1 in x 5 inx 17 in for DWTT specimens versus 1 in x 10 in x 20 in forMIG weldments). Tests of base plate material from otherheat treated MIG weldments (Figs. 16 through 19) did notshow the same discrepancy compared to is-received baseplate values (5) or to MIG weliment values (Table 9),the relation between specimen sizes being the same.

In other base-plate tests, high interstitial unalloyedtitanium was shown to be sensitive to SCC (Kiscc/KIx = 40/69)while all other alloys tested were relatively uiaff-!ctedby the salt water environments.

The -es 1 lts of tests conducted on MIG weldments both,s-welded and solution heat treated are sho' n in Table 9and in Figs. 28 through 34. Base plate values for bothconditions are shown for comparison. The procedures usedin fabricating the weliments and the mechanical propertieso .)f the •eldments are des,'ibcd in detail in reference 7.

The procedure in SCC testing the wei.-nents has beendescribed in the saim reference- As in dicated in Table 9,-pecimen,, for tests at Aeld centerline () , at the heataffected Zone (HAZ) ( 3 "`16-in), and in the as-receivedba•se plate fo, al loys T94 and T95 we-v ' /2-in. *ide withI "32-in. deep side-grooves. Sp.ci-.n dimensions for otherplates ire irdica'ici ir the table.

The values of Ki•c-c:. for a. -received base platematerial were higher than "h-,-e for the ield depi-sit andUIAZ. both as-&elded and h-<,at treated. Alsu, the K scc valuesfo.- heat !reated plate rrterial generally exceeded the Klsccvalues folr the heat treated *eid metal eposits and HAZ.Va'lues of Kls(.c uere generally higher at "t-td centerline

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than at the HAZ for as welded plate, while the opposite wastrue for the heat treated weldment. Heat treated weldmentshad generally higher Kiscc values for both weld centerlineand HAZ compared to as deposited weldments.

FRACTOGRAPHIC STUDY OF FATIGUE FRACTURES

In a recent low cycle fatigue study of Ti-7A1-2Cb-lTain air and in salt water environments, electron fractographymethods revealed tlie presence of quasi-cleavage fracture onthe fracture surfaces of the specimens tested in salt water(15). Quasi-cleavage fracture, which indicates SCC was alsoobserved at the fatigue-fast fracture interface on a speci-men tested entirely in air. Since a high likelihood existsthat such fractures can be caused by moist air in alloyswhich are highly sensitive to SCC, a number of single-edge-notch (SEN) fracture mechanics specimens of titanium alloyswere examined for evidence of quasi-cleavage fracture in thefatigue cracked portion of the fracture surface and at thefatigue cr3ck-fast fracture interface.

The eight specimens examined by electron fractographicmethods were selected to represent the variety --f Kic, yieldstrength (YS), and drop-weight tear energy levels availabl;.Preparation of the SEN specimens had included fatiguing themin three point bending at approximately 30% of YS to generatea fatigue crack before they were fractured in tension. Table10 shows the composition and Kic values for these alloys alone,with SCC test parameters (Kix and K ). The SEN specimenswere tested in the heat treated condion shown, while SCCvalues were taken from mill-annealed material, with theexception of T21, which was SCC tested in the heat-treatedcondition shown.

The fractographic method used was the standard two stepmethod of replication described in reference (15), withexamination of the replica in the electron microscope atrelatively low levels of magnification (3000X-6000y).

Cleavage or quasi-cleavage was observed on some of thespr-cimens, while others spowed none (See Table 10). Thepresence of this low energy fracture mode indicated that some kcrack growth, especially in the fatigued region, was due toSCC. In the specimens where cleavage or quasi-cleavage wasevident the fatigue-fast fracture interface could not beclearly defined, therefore, whether this mode of fractureextended into the region of the fracture surface associated V.A

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LOW CYCLE FATIGUE CRACKPROPAGATION IN A2OlB, A302B, and A517F PRESSURE VESSEL STEELS

(T W. Cr7Jker and-JA7Lng-

This report describes the results of an investigation oflow cycle fatigue crack propagation in three structuralsteels widely used in welded pressure vessels: A2OlB, A302B,and A517F. The purpose of such a study stems from the factthat the occurrence of small sharp defects in large weldedstructures cannot be totally eliminated. Careful fabricationand nond(,structive testing procedures are not sufficient tocompletely guarantee the absence of small fabrication defects.Therefore, low cycle (i.e., failure in less than 100,000cycles) fatigue crack propagation is a potential failure modefor large welded structures which undergo repeated loadingexcursions.

The crack propagatiop phase of fatigue failur.? is arelatively new area of study. It departs from the tradi-tional design approach to fatigue which combines crackinitiation plus crack propagation into an arbitrary defini-tion of fatigue failure. Ho,',ever, fatigue crack propagationstudies are a promising means -of providing a more rationalapproach to structural fatigue situations Ahere exper--ii-ii-,suggests a very high probability if .. r- - c:xsting fia1s.

Attention is currently limited to pl;,te materials fortwo reasons: the inhierent s-mplicity of plate materialsrather than weldment:a for an initi.il research program andthe fact that. although %Ided joints, are a pr•ime source ofdefects, the propagation ;-, defects in fatigue is controlle.by the stress pattern an,' is by no means limited to weld metal.

The steels chosen fc.r this investigation are importantbecause of their broad application, but are of particularsignificance in this studly be(:c.u,-e the fatigue test speci-me n emplo)yed] Aere taken rtm the shell mater ias of full-si e pressure vessels. The pressure vessels themselve:s hadfirst undergone fatigue te.--ting as part of the WeldingResearch Council'.s Pressure Ve- sel Rcscart-h Commiu-ttee (PVRCIprog-ram of testing full-:i:e pressure vessels. The finalfailur" anallyses conducted on tbf six PNRC vessels are pub-lished in referenc,:s 16-19. The subseqjent laboratory

59

Fig. 35 -Quasi-cleavage facets in fatigued zone of unalloyed titanium(T-17) SEN specimen. Note fatigue striations almost completelysur-rounding facets. 3000X

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Fig. 36 -Cleavage fracture in Ti-8A1-lMo-lIV alloy(T-19) SEN specimen. 300OX

61

Fig. 37 -Low energy fracture in Ti-6A1-4Zr-ZMo alloy

(T-55) SEN specimen. 3000X

62

investigation discussed in thi' report describes the lowcycle fatigue crack propagation characteristics of thesesteels in air and in two aqueous environments, simulatedboiler water and 3.5% salt water. In addition, approxi-mate correlations between the results of these laboratorytests and the actual vessel fatigue lives are shown.

EXPERIMENTAL

(a) TEST SPECIMEN AND MATERIALS

Plate bend fatigue specimens, Figure 38, were machinedfrom 2-inch thick vessel shell materials of A2OlB, A302B,and A517F steels. The A2OIB and A302B steels were normal-ized and the A517F steel was quenched and tempered. Chemicalcompositions and tensile properties of these steels aregiven in Tables 11 and 12, respectively.

(b) APPARATUS AND PROCEDURE

The experimental data reported herein are based on theoptically observed macroscopic growth of fatigue cracksacross the surface of center-notched plate bend specimens.The specimens are cantilever loaded in full-reverse (tension-to-compression) strain cycling. These procedures are treatedin detail in reference 20 and their development and priorapplication are contained in previous reports in this series.

An engraved center surface notch is placed at theminimum thickness of the test section to assure rapid fatiguecrack initiation at relatively low cyclic strain levels. Afoil-type resistance strain gage is placed on the test sec-tion well ahead of the advancing fatigue crack to monitornominal surface strains, which are recorded as total strainrange values. The progress of the fatigue crack across thetest section is observed from directly above through a slidemounted optical micrometer. Extensions at both ends of thecrack are monitored and the sum is reported. Automaticdeflection control of the test is obtained through micrometeradjusted microswitches which contact the specimen arm at themaximum and minimum points in the loading cycle. Specimensare cycled at a rate of approximately five cycles per minutein a sinusoidal loading pattern.

Based on extensive experience with this test method,it has been found that, in the plate bend specimen, constanttotal strain range cycling results in a constant value of

63

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66

lfAtigue crack growth rate. For this reason, all tests reportedin this paper were conducted under constant total strain rangeloading. Such loading is obtained by gradually reducing thedeflection control limits as the fatigue crack increas.b ilnsize.

Cycling is begun at the lowest desired strain range andis continued until a fatigue crack is initiated, propagatesaway from the mechanical notch, and establishes a constantrate of growth. Then the specimen is loaded to the next higherstrain rdnge level desired and cycled until the new, higherrate of growth is established. Using this procedure, onespecimen can be used to generate several successive datumpoints.As a general practice, data on a material are first taken inzir and then the test procedures are repeated with fresh speci-mens under the aqueous environments.

The corrosion cell employed in the aqueous environmenttests is placed over the portion of the test section contain-ing the fatigae crack. It is made of molded polyurethanewhich is relatively soft and flexes with the specimen. Theaqueous solution is constantly circulated through the corrosioncell from a reservoir by a small electric pump. The solutionundergoes filtering and periodic monitoring of its pH valueand salt content. A glass cover at the top of the corrosioncell permits optical observation of the fatigue crack.Periodic removal of tarnish or rust from the test surface isrequired to maintain good visibility of the crack.

Wet fatigue studies were conducted with two environments,simulated boiler water and 3.5% salt water. The simulatedboiler water is the same comPoisition used for internal pres-surization of the PVRC full-size pressure vessels duringfatigue testing. Therefore, its effect on the fatigueperformance of the three pressure vessel materials discussedIn this paper is of special interest.

The chemical composition and preparation procedure ofthe simulated boiler water is as follows: 4.0 gramz of sodium 4-

hydroxide (NaOH), 7.6 grams of sodium sulfite (N92 S0 3 ), and0.09 grams of copper sulfate (CUSO5 were each dissolvedseparately into two liters of dist I water, forming threesolutions. Fifty ml from each of these three solutions wasthen added to one liter of distilled water to produce thefinal mixture. A pH value of 10.5 to 11.0 is obtained fromthis mixture. Tannin was included in the PVRC simulated

67

boiler water but was eliminated from the NRL solution becauseit interfered with optical observation of the fatigue crack.

RESULTS AND DISCUSSION

(a) FATICUE CRACK PROPAGATION IN AIR

Fatigue crack growth rate data from the air environmenttests are plotted versus total strain range on log-log coordi-nates in Fig. 39. The data for A201B and A302B steels fallon a common third power (3:1) curve and the data for A517Fsteel fall on a fourth power (4:1) curve. The equationsbetween fatigue crack growth rate and total strain rangewhich are established by these plots h? °e the form

d(2a) c(C ) m&N -

d(2a)where -T_ is the fatigue crack growth rate, CT is the totalstrain range, m is a numerical exponent (slope), and c is anumerical constant (intercept). These fatigue crack growthrate versus cotal strain range plots provide baseline rela-tic nshipsthat illustrate the intrinsic resistance of these-at.2rials to the growth of fatigue cracks. This informationis used fcr comparing the fatigue crack propagation charac-teristic& ot materials, determining the effects of aqueousenvironmnnt., on fatigue crack propagation and estimatingstructuril fatigue life.

Consider first the sensitivity of the three steels tocyclic strain. The data in Fig. 39 indicate that, on astrain range basis for load intensity, the two normalizedsteels A201B and A302B display very similar low cycle fatiguecrack propagation characteristics. The quenched and temperedA517F steel displays distinctly different characteristics.Previous testing experience suggests that the locus of thedata on Fig. 39 for Lhe normalized steels is not shared byany other material, i.e., no relationships with slopes as lowas three have been observed among a wide variety of ferrousand nonferrous metals previously tested by these methods.However, the locuE -• the data for the quenched and temperedsteel is very c. that of other high strength, quenchedand tempered ste-. ch as HY-80 and 5Ni-Cr-Mo-V (20-24).

A second compirative perspective is shown in Fig. 40.This is a summary log-log plot of the air environment fatigue

68

1000 T I I 1 1000

PRESSURE VESSEL STEELS 0

[ FULL-REVERSE STRAIN CYCLEAIR ENVIRONMENT

I -J

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range for A•01 ci, A30 28,an A517F t e nan m i nt ir nv on

ment. The Lndex of Struct rral Fatigue Life is indicated along the

right-hand ordinate.

69

t t

1 I 11111I:100

- PLATE BEND FATIGUE TEST '"

FULL-REVERSE STRAIN I "CYCLE

AIR ENVIRONMENT ,

P12Sj

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II

11 • 1' i_ • II 1,000,0000.1 0.2 0,4 0.6 .0 2.0 4.0 6-0 00

RATIO OF TOTAL STRAIN RANGE TO PROPORTIONAL LIMiT STRAIN RANGE

Fig. 40 - Air environment low cycle fatigue crack propagation charac-

teristics of AZ01 B, A302B, and A517F steele on a n o r maIi z ed load

intensity baslis I

70

crack growth rate versus the ratio of total strain range toproportional limit strain range for the three steels con-sidered in this report. The proportional limit of each steelin the plate bend specimen configuration is indicated on thetotal strain range axis of Fig. 39. Using this referencepoint, a normalized load intensity factor is obtained on Fig.40 which is comparable to yield strength loading at a f&rctorvalue of unity.

Looking at Fig. 40. it can be seen that when fatiguecrack growth rates for the three steels are compared on thisnormalized load intensity basis, a different order appearsthan what is suggested by the curves in Fig. 39.

Significant differences exist among these materials withregard to the rate of fatigue crack propagation at total strainrange levels which are in fixed relation to thte proportionallimit strain range. On this basis, it can be seen that A201Bpropagates fatigue crp'!s less rapidly than either A302B orA517F, which display similar characteristics in this respectBy the same reasoning it can be shown that A2UJiB possesses agreater tolerance for plastic strain under cyclic loadingthan either A302B or A517F. For example, Fig. 40 shows thatat a total strain range equal to the proportional limit ineach material, fatigue cracks propagate approximately fourtimes faster in A302B than in A201B and fi,'e times fasterin A517F. The higher fatigue crack growth rate in the steelswith higher yield strength are more largely due to increasedloading (strain range) rather than to a decrease in theintrinsic resistance to the growth of fatiguc cracks.

This discussion should not be interpreted as a condemna-tion of higher strength materials. On the contrary, it merelypoints out little understood but important differences in thefatigue crack propagation characteristics of metals at variousstrength levels and provides a guide for their safe use inhigh performance structures. Since structural design andfabrication processes can alter the relative merit of aspecific material, these aspects, as perceived from the PVRCfull-size pressure vessel tests, will receive further con-sideration in a later section of this report.

(b) FATIGUE CRACK PROPAGATIOY IN AQUEOUS E.NVIRONMENTS

The results of the fatigue crack propagation tests con-ducted in the two aqueous environments, simulated boiler water -

71

and 3.5% salt water, are shown in Figs. 41 through 44. ThesefigureS are log-log plots of fatigue crack growth rate versustotal strain range. In each case the air euvironmentplot for the corresponding material is indicated by thedashed line. The effect of the aqueous environment isillustrated by the vertical displacement of the aqueousenvironment data to higher crack growth rates than wereobserved in air environment tests at corresponding valuesof total strain range. For convenience, data for the A201Band A302B steels are plotted on common curves, Figs. 41 and42, and the A517F data are plotted separately, Figs. 43 and44,

As might be expected, fatigue crack propagation in eachof the three steels was generally adversely affected by thepresence of an aqueous environment. Exceptions to thistrend occurred at low total strain range levels in the A201Band A302B -teels, where the environment appeared to have littleor no effect. All three steels were affected more severelyby the 3.5% salt water than by the simulated boiler water.

Generally. the effect of the aqueous environmentsinciroased in proportion to the total strain range level.The salt water environment data for all three steels show suchabrupt transitions as to suggest the existence of a thresholdlevel of cyclic strain below which the materials are rela-tively insensitive to the environment, in a manner similarto the Kiscc concept for stress-corrosion-cracking (12). Alsothe salt water environment data for each of the three steelsappeared to indicate a maximum environmental affect in themid-range of total strain range values, with the environmentaleffects diminishing somewhat at peak total strain range levels.However, these tests were conducted under 5 cycle per minutesinusoidal loading with no hold period at maximum load. There-fore, loading rate effects may be reflected in these data andthis must be considered when the information is applied tofuil-size structural situations involving slower rates ofcycling.

inally, experience suggests that an order of magnitudeincrease in fatigue crack growth rates due to the introductionof an aqueous envircnment provides a first-order benchmark forjudging the performance of structural materials. Most highstrength structural steels exceed this benchmark and someultrahigh strength stee>-s grossly exceed it (21). Lookingagain at Figs. 41 through 44, it can be seen that the maximum

72

GO000 - ,1r---r rn , T--11000

X1A2OI AND A3028 SI EELSFULL- REVERSE STRAIN CYCLESIMUJLATED BOILER WATER ENVIRONMENT

/ ---- PLOT OF A20IB AND -J"I A302 STEEL

DATA !N AIR

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4 A A302B STEEL 58 KSI Z

; PROPORTIONAL LIMIT STRAIN RANGE(500 MICROINCHES/INCH PLASTIC STRAIN RANGE)

A2018 A3026

1.0 1 ý I I I I _. I I I II - - '1,00 0 ,0 0 0

1000 10,000 100,000

TOTAL STRAIN RANGE (MICROINCHES/INCH)

Fig. 41 - Log-log plot of fatigue crack growth rate versus totalstrain range for A201B and A30ZB steels in a simulated boilerwater environment

73

00,I 11 1 I I I I I II 000

A201B AND A3025 STEELSFULL-REVERSE STRAIN CYCLE3.5% SALT WATER ENVIRONMENT

a:/0/ "---PLOT OF A201B AND _

I. A302B STEEL

DATA IN AIR oLU IOH-/• OOOu

m - U Uz-'r

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l_ A302B STEEL 58 KSI

PROPORTIONAL LIMIT STRAIN RANGE(500 MICROINCHES/INCHPLASI IC STRAIN RANGE)

A201B A302B

1.0 I I I I I I , 1,000,0001000 10,000 100,000

TOTAL STRAIN RANGE (MICROINCHES/INCH)

Tig. 42 - Log-log plot of fatigue crack growth rate versus total strainrange for AZ01B and A302B steels in a 3.5% salt water environment

74

1t00 I I I I I 1000

A517F STEELFULL-REVERSE STRAIN CYCLESIMULATED BOILER WATER ENVIRONMENT

LUJ/Ww -..J OF A517F STEEL

DATA IN AIR U.I 0I .-

w o

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SPROPORT04AL LIMIT STRAIN RAkNGE1(500 MIOROINCiHES ANCH PLSI STRAIN RANGE

A51TF

1.0 1 .1 1, i I I ! I I I J , J I , I I M100 0 . 0 0I000 10,000 100,000

TOTAL STRAIN RANGE (M IC R01NC I IE-/ INCH)

Fig. 43 - Log-log plot of fatigue crack growth rate versus total strainrange for A517F steel in simulated boiler water environment

75

S I

1000 1 ~~000

- A517F STEELFULL-REVERSE STRAIN CYCLE

_ 3.5% SALT WATER ENVIRONMENT

w

.J h-PLOT OF A517F STEEL 7%IL.

U. I DATA IN AIR

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(500 M11CROINCHES/INCH PLASTIC STRAIN RANCE)

A517F

I.O 1,000,0001000 10,000 100,000

TOTAL STRAIN RANGE (MICROINCHES/INCH)

Fig. 44 - Log-log plot of fatigue crack growth rate versus total strainrange for A517F steel in a 3.5%1o salt water environment

76

9.

increase in fatigue crack growth rates due to the aqueousenvironments in the A2OlB and A302B steels lie well withinan order of magnitude for the simulated boiler water andthe 3.5% salt water environments. The A517F steel dataexceed an order of magnitude increase in the salt waterenvironment but fall well below this increase in the simu- Ilated boiler water environment.

(c) COMPARISONS WITH FULL-SIZE VESSEL TEST RESULTS

The results of this investigatton present a rare oppor-tunity to compare laboratory fatigue crack propagation datawith full-size structural tests. I.w cycle fatigue crackpropagation is gaining recognition as a critical failuremode in cyclically-loaded large welded structures. Inparticular, there is general agreement that fatigue crackpropagation was the dominant fatigue failure mode in thePVRC vessels. However, scant data are available to enablethe establishment of proper criteria for designing againstthis mode of failure. As a first step toward this goal, theauthors have proposed the Index of Structural Fatigue Lifeas an initial attempt to correlate the effects of cyclicstrain, mean strain, and aqueous environments with thefatigue life of structures (20).

The Index of Structural Fatigue Life is defined as thenumber of cycles of repeated load required to propagate asmall subcritical flaw to 1-inch length at a given intensityof loading, as measured by the total strain range, in theplate bend fatigue test. This Index is shown along theright hand ordinate of each data plot, Figs. 39 through 44,Although it is an arbitrary failure criterion, it servesas a helpful means of translating crack growth rate datainto a more commonly recogniz'3J form of fatigue data.

Figure 45 shows a log-log piot of total strain rangeversus cycles to failure for the PVRC vessels (19). Threedata symbols are indicated for each of the six vessels: theactual vessel failure based on the measured or estimatedtotal strain range at the failure location, a predictedfailure point based on the Index of Structural Fatigue Lifefor the vessel shell material in a simulated boiler waterenvironment, and a predicted failure point based on theIndex of Structural Fatigue Life for the vessel shellmaterial in an air environment with a correction for theeffect of mean strain. The mean strain correction consistedsimply of taking the fatigue crack growth rate for full-reverse strain cycling, Fig. 39. and multiplying it by a

77

0

X 0U)w ww a

0,.• W.0Z2

w W x '544o

, ,, %L L 9 0, 0

,-'n".-• - .o o

-- %1 -0.> ._I 1

• , O, 0"~~~~~ t:) So.:- .t

IID

0a 0

wU

a' 0

d -d

0

0 0 00

- ~(H)NIiS]HDNIO•JIbNW) 39NV• NIV•1S iviOi

Fig. 45 - Log-log plot of total strain range versus cycles to failure forthe PVRC full-size pressure vessel test results. Actual failure pointsplus failure predictions based on the Index of Structural Fatigue Lifeare indicated.

78

factor of three to obtain the approximate Index of StructuralFatigue Life value. Laboratory investigations have shownthat a threefold increase in fatigue crack growth rates isthe minimum affect of mean strain on structural steelscycled to strain amplitudes near or into the plastic strainrange (20, 22).

The comparison of actu2l failurez, versus lauoratorypredictions is rather encouraging for an initial effort. Thepredictions for vessels 1 (A201B) and 3 (A302), which failedin the cyclic life span between 7000 and 9000 cycles, agreedvery closely with the actual life to failure. For vessels2 (A2OlB) and 4 (A302B), which failed in the life spanbetween 40,000 and 90,000 cycles, the predictions were con-servative. The closest agreement for vessels 2 and 4 camefrom predictions based on the simulated boiler water envir-onment.

The opposite relationship between predicted and actualfailure points was seen for vessels 5 and 6 which were madeof A517F steel. Actual cyclic life to failure in thesevessels fell well below the predicted failure points. Forthe A517F steel, the predicted fatigue lives for each of thetwo criteria considered, simulated boiler water environmentor air environment corrected for mean strain were identical.

Based on fatigue life at a specific total strain range,the largest factor separating actual and predicted failureswas 2.4. Considering the inexactness of fatigue testing ingeneral and specifically the fatigue testing of large weldedstructures, this represents a good first approximation.Further experimentation and refinements upon the crack propa-gation approach to fatigue failure would seem to offer a morerealistic method of predicting failures for the "worst case,"i.e., the case of the large welded structure which containsundetected sharp flaws which preclude a structural lifeperiod for crack initiation.

An additional aspctt of the fatigue life data in fig. 45illustrates a potential advantage in the use of higher strengthmaterials. Vessels 1, 3, and 6 were constructed to identicaldimensions and cycled at the same maximum internal pressuriza-tio)n (4400 psig). However, vessel 1 was made of A201B steel,vessel 3 was made of A302B steel, and vessel 6 was made ofA517F steel having yield strengths of 48. 58, and 107 ksirespectively. Refer-ing back to Fig. 45, we see that with

79

higher yield strength, plastic strains were minimized.Consequently, a smaller total strain range existed at theconstant value of internal pressure, resulting in longerfatigue life. Specific life values were 7,516; 8,500 and19,272 cycles for the A201B, A302B and the A517F vesselsrespectively. This series of experiments serves to illus-trate one manner in which the potential advantages ofhigher strength materials can be exploited while minimizingthe potential disadvantages of such materials. The inter-action of the material, structural, and eiiv±ronmentalparameters are highly complex. Additional precise structuraltests are required to develop a comprehensive fatigue lifemodel incorporating all of the parameters affecting thegrowth rate of fatigue cracks including mean strain, stressstate, and effect ot flaw size on strain range.

SUMMARY AND CXONCLUSIONS

Low cycle fatigue crack propagation studies were con-ducted on A201B, A302B, and A517F pressure vessel steelsin ambient air and in two aqueous environments, simulatedboiler water and 3.5% salt water. The study employedplate bend fatigue specimens which had been machined fromthe shell materials of full-size pressure vessels. Thevessels had previously undergone fatigue testing under theauspices of the Pressure Vessel Research Committee.Empirical correlations were made between the fatigue crackgrowth rate and total strain range. These data were thencompared to the fatigue life of the six PVRC vessels.

The following conclusions have been reached from thisstudy:

1. Fatigue crack growth r-ates in A201B and A302B steelsin an air environment follow a common third power relation-ship with total strain range. For A517F steel in an airenvironment a fourth power relationship with total strainrange was observed. The locus of the relationship for A2OIBand A302B is unique among structural steels tested to date:whereas the relationship for A517F is very similar to thatobserved for other high strength, quenched and tempered steelssuch as HY-80 and 5Ni-Cr-Mo-V.

2. Comparisons of fatigue crack propagation characteristicsamong the three steels tested, based oTn total strain rangeunits normalized with respect to the respective proportional

80

limits, indicate that A2OlB possesses the greatest tolerancefor cyclic plastic strain. A302B and A517F exhibit similarbehavior in this regard. At a total strain range equal tothe proportional limit in each material, fatigue crackspropagate approximately four times faster in A302B than inA201B and five times faster in A517F.

3. Fatigue crack propagation in all three steels wasadversely affected by the presence of an aqueous environment.The 3.5% salt water environment proved to be more severe thanthe simulated boiler water environment. Generally A517F wasmore sensitive to environment than A2OB or A302B, althoughnone of the three steels exhibited a gross degradation infatigue properties due to an aqueous environment.

4. Comparisons between an Index of Structural FatigueLife based on these fatigue crack propagation data andresults of the PVRC full-size vessel tests are reasonablyclose for the A2OB and A302B steels. Less satisfactoryagreement was found for the A517F steel in that the pre-dicted .alues were not conservative.

CORRELATION OF TWO FRACTURE TOUGHNESS TESTS

FOR RT-H-STRENGTH STEELS

(C. N. Freed)

Linear elastic fracture mechanics allows measurement ofthe rate at which strain energy is released upon crack exten-sion from a sharp flaw or fatigue crack. The stress intensityfactor, K ic which may be calculated is a property of thematerial. As the test is based upon the assumption thatplane strain exists in the center of the specimen, thematerial to be tested can not possess high toughness, i.e.,it must fracture under conditions of elastic loading. Thechief advantage of this fracture toughness is that KIc is afunction of the length of the crack and of the stress actingupon the crack; thus, if any two of these values are known,the third may )e calculated.

A preliminary correlation has been developed between theto fracture toughness parameters, KiC and DWTT, for a varietyof steel alhoys. This correlation permits predictions to benzide of the expected range of Kic from drop-weight tear test(DWTT) energy values and reasonable estimates of flaw sizeand stress level for steels in which a crack %ill propagatewithout gross plastic deformation.

81

MATERIALS AND PROCEDURE

The steel alloys which were investigated include 12-and 18-Ni maraging steels, 9-4-0.25 (Ni-Co-C) quenched-and-tempered, 4140, D6-AC, and 5Ni-Cr-Mo-V steel. (Table I)About one-half of the specimens were tested in the "as-received" mill condition while the remainder were heattreated at NRL as part of a program headed by Dr. P. Puzak.The specific mill processing and NRL heat-treatments havebeen reported in previous Quarterly Reports (5, 24, 25)The mechanical properties presented in Table 13 are alsodue to Dr. Puzak.

The single-edge-notched specimen (SEN) was used todetermine the KIc values reported in Table 14. The dimen-sions of this specimen are similar to those reported in theSe-enth Quarterly Report (25). In both the DWTT and theKic investigations, the specimens were approximately one-inch thick (except J-79); this usually represented thefull-plate thickness although in several cases the specimenswere cut from 2-in.-thick plate. Many of the Kic specimenswere side-grooved to a depth of 5 percent of the thicknesson each side. A fatigue crack of 0.10 in. was formAed atthe tip of the edge notch.

Load-displacement graphs for each KI, specimen weredrawn by an X-Y recorder. When initial aeviation fromlinearity occurred at or very near maximum load, thisvalue was used to calculate Kic; otherwise, the load atthe lowest, distinct instability was chosen for the calcu-lation. The detection of the load instability was madewith a beam displacement gage in-;tramented with a straingage circuit. A plastic zone correction (one iterati on)was included in the determination of Kic; however, Table 14also indicates the KIC value before inclusion of the plasticzone correction.

ANALYSIS OF DATA

In Fig. 46, Kic is plotted ag:ainzt yield ;•rength (YS).The graph indicates that ;tn inver.-se r-elatti.nsrip exist- betweenKjc fracture tough-ness and YS for high strength steels. Itis evident that a ';ide r:inrge of KIC values can exist ait .i givenYS for various alloy :-nteels ari, he.:at tr<-.tmeints. The fracturedirection is indicated as RW (hortzont~A lines) and WR (vterticallines) according to ref, 14. Each pý:int represe'.s an average

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83

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HIGH STRENGTH STEELS280__LEGEND

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- 20-0 1 -0i t I I I I, I II .

140 160 180 200 220 240 260YIELD STRESS (ksi)

Fig. 46 - KIc plotted against yield strength for a varietyof high strength steels

85 *

of the Kic values obtained from that particular alloy and heattreatment.

The plot of K versus DWTT energy is found in Figure 47.Although some scatier is present, it is evident that a directcorrelation exists; an increase in KIc values correspondsdirectly with an increase in DWTT energy. The vertical linedrawn at 3500 ft-lb on the DWTT energy scale represents theapproximate value at which the fracture surface of the DWTTspecimen becomes 100 percent slant fractilre. The line servesa second purpose in that at this point the nominal-to-yield-stress ratio in the K test approaches 1.0. Beyond thisvalue, the crack may A extending under conditions of grossyielding rather than by linear elastic fracture mechanics.

There are several factors which might affect theaccuracy of this preliminary correlation. Since KIc is afunction of strain rate, the high rate of applied strain inthe DWTT may cause the specimen to fracture with a lowertoughness indication than would be expected if the strainrate:, of the two tests were similar. However, as the steelstested are relatively insensitive to strain rate effects,this factor should not be too significant.

An energy loss is inherent in the fracture of thebrittle, crack-starting weld. Altnougn tne loss is small,it might be a significant portion of the recorded energy oftne most brittle DWTT specimens. It would be expected thatthis factor should cause the curve to intersect the abscissaat some finite value of DWTT energy rather than zero DWTTenergy. This may be demonstrated when more data are obtained.

A third and perhaps most important point is that as thecrack tears through the DWTT specimen, it may initially begoverned by a plane strain state of stresses but in thetougher materials it will propagate primarily under mixedmode conditions. The mixed mode is probably caused by thelateral expansion of the crack at the same time that itmoves forward. The lateral movement would effectivelydecrease the constraint around the crack tip and cause thestress intensity factor to rise since plane-strain conditionsno longer exist. Thus, the plastic zone would increase inthis region of mixed mode fracture, and this would eventuallyproduce surface relaxation evidenced by shear lips. The resultis that the energy measured in the DWTT would. for toughmaterials, represent mixed mode while the stress intensity

86

HIGH STRENGTH STEELS, 280-

' 260 LEGENDSYMBOL CHEMISTRY

240 12 Ni MARAGE* l18 Ni MARAGE

v 220 - 9-4-0.250 *W 4140 , *

(C2 200- RW0 4WRI 4 z-•180-

C) 160 "V)

"' 140 ,

120 b

SI00-

80 +

~60-

W 40--J

• 20--O0 I I I I I 8

800 1600 2400 3200 4000 4800DWTT ENERGY (ft-lb)

Fig. 47 - Preliminary correlation of K withdrop-weight tear test energy Ic

87 4

factor determined in the Kic test uoLld apply to plane-strainconditions. This discrepancy may result in an increase ofthe correlation scatter at higher DWTT energy values.

The DWTT energy is plotted against fIc in Figure 48. Inboth tests the resistance of the metal to crack propagationis a function of the plastic zone size, and OIc is an indica-tion of this factor for the KIC investigation. The normali-zation of KIc by the yield stress may reduce the degree ofscatter.

Ii Figure 49, the strain energy release rate,hic, iscompared to the DWTT value divided by the nominal fracturearea (excluding the brittle weld). This plot was drawn inorder to compare similar quantities on each axis (in-lb/in2 ).The least amount of scatter in any of the figures is indi-cated in this curve.

In conclusion, it should be emphasized that these arepreliminary correlations based on a limited amount of data.As more information is accummulated and the degree of scatterthat can be expected is better defined, the usefulness of thesecorrelations in permitting predictions of KIc from DWTT energyvalues will become better established.

ALUMINUM ALLOYS(R.W. Judy, Jr. and R.J. Goode)

*STRESS CORROSION CRACKING

Exploratory stress-corrosion-cracking (SCC) studies have beenconducted on aluminum alloys in a manner similar to that usedfor titanium alloys, which is described in an earlier sectionof this report. Previous studies (5, 7) have shown that thesemethods do not yield meaningful results when applied toaluminum alloys except for specimens oriented in the "TW" or"TR" (14) position; therefore, these studies of aluminum alloyshave been confined to the "TW" or short transverse orienta-tions. Specimens of the proper orientation were obtainedfrom one inch plate by welding tabs to a 1 in. x 1 in. barand cutting the 1 in. x 1/2 in. x 1 in. specimens from thewelded assembly.

Stress-corrosion-cracking data and some mechanicalproperties of three alloys tested are shown in Table 15 andFigures 50 through 52. Two of the alloys, 7079-T6 (A13) and7075-T7351 (Ai4) were relatively sensitive to SCC, havingKIscc values of 11 and KIx values of 22.8 and 25.3 ksi . in,respectively. The 7005-T63 was insensitive with a Klsccvalue of 27.5 and Kix of 27.5 ksi .in. The tests were con-tinued over a long time period in a stagnant 3.5% salt water

*This research was supported by the Advanced Research Projects Agencyas part of the ARPA Coupling Program on Stress-Corrosion Cracking.

88

HIGH STRENGTH STEELS

2.00- S LEGENDSYMBOL CHEMISTRY

12 Ni MARAGE18 Ni MARAGE

1.80 a 9-4-0.25* 14140

4*RW

1.60 4WR

1.40

1.20

1.00 4

0.80-

0.60 -4

0.40 -

0.20

0 800 1600 2400 3200 4000 4800DWTT ENERGY (fi-lb)

Fig. 48 - Preliminary correlation of e 1¢ withdrop-weight tear test energy

89 '3

HIGH STRENGTH STEELSLEGEND

2400 -SYMBOL CHEMISTRY

A 12 N1 MARAGE* 18 Ni MARAGE

2100 - 9-4-0250 4140

*RW

1800 *WR

CQ1,500-

S1200-

900-

600-

300-,

0 1000 2000 3000 4000 5000 6000DWTT ENERGY (ft-lb)

I .- ' --I I

0 4,000 8,000 12,000 16,000 20,000 24,000DWTT ENERGY/AREA (in-lb/i092)

Fig. 49 - Preliminary correlation of strain energy release rate,

Ic with drop-weight tear test energy resultsg

90

*-4J N N N

- p4

004

0~0

.1c x o0- 4.)

VI~4. 00

* 5

r-44 0-

000

- 54 4.X

V cn to

0 0

t- 0

o 0 0

t--

Cl)91

N 0

z z

-0

z 001.

(A 0

00 u 0

0 0-4

w 0tU0 U

u

0)

-0-

10 0 U' 0kiN~ N ~ -

CNIl' IS)4) IN4

92

0

z0n

0

If) - W~

1-n~ 00I- 0

0u~ 4

III 0x3 0

0ý u

-0u

uuV)

004

qy-m

93~

V z-- 0

rr 03.

U,'

I (.4cr0 a

0iL

944

solution which was changed twice daily, possibly affectingthe absolute Kis c values obtained; however, the generalindication of reiative sensitivity or insensitivity isconsideree to be valid.

FRACTURE TOUGHNESS INDEX DIAGRAM FOR ALUMINUM

The fracture toughness index diagram (FTID) for alumi-num is shown in Fig. 53. No additions or modificationsof these FTID have been made for some time including theperiod covered by this report. It is being presentednow for the benefit of those who have recently becomeacquainted with this report series since it was not coveredin the Eleventh Quarterly. The aluminum FTID reflects thecorrelation of drop-weight tear test (DWTT) and explosiontear test (ETT) performance with yield strength (YS). Theoptimum materials trend line shows the limiting or optimumvalues in the weak fracture direction with respect to theprincipal rolling direction in commercially provided plate.Correlations of ETT performance with DWTT energy valuesare shown on the left of the diagram. Alloys with DWTTenergies that fall below the elastic-plastic transitionband (-300-400 ft-lb DWTT e-orgy) can be expected to supportrapid crack propagation (catastrophic fracturing) at elasticstress levels. Increasing DWTT energies above this transi--tion band indicate increasing plastic strain capabilitiesbefore a crack will propagate. Above 60 ksi YS all alloyswould be expected to be "brittle" anr fracture mechanicstechniques will be required to characterize fracture tough-ness. Below about 50 ksi YS a wide variety of alloysshould be available that require extreme amounts of plasticstrain for fracture.

HIGH STRENGTH STEELS(P. P. Puzak)

The fracture toughness characLeristics of steels andwelds of approximately 130 to 200 ksi yield strength havebeen studied. In this report are the results obtained for3/4 to 4-in. thick plates of a 5Ni-Cr-Mo-V steel and 1-in.thick plates of a 9Ni-4Co-0.2C (9-4-0.2C) quenched-and-tempered (Q&T) alloy steel. Fracture toughness investi-gations were also conducted f,'r pro'ucer fabricated manual(STICK) and metal-arc inert gas (MIG) welds of the nominal2Mn-2Ni alloy composition developed for welding the5Ni-Cr-Mo-V c-teel, and NRL fabricated, fully-automatic,tungsten-arc inert gas (TIG) %elds of the 9Ni-4Co-0.2Calloy and 12Ni-3Cr-3Mo and l7Ni-2Co-3Mo maraging steelweld metal compositions.

95

ALUMINUM ALLOYS

ONE- iNCH PLATE AT 30-F

Fu)3000 0" q~RI ALL',"

OPTIMUM MATERIALS 0 0 5086-HI12U- TREND LINE 0 5083-0

E- ] \ WEAK DIRECTION" A ,H>" 15 % + STRAIN I, 5456-H321

1J t 2024-T4.z2O00 -- ", C•0 2219-T87

a '7075-T6D 7079-T6

"8-15% STRAIN > 0 1106-T63

- x X 7005-T63H 7075-T7351

- 1000 > 2219-T851- 3-•8% STR"'N

0

% T STRAIN

GENEALLY FLAT BREAKE .- _XTENSIVE FRCTUR _ _ _

0 10 20 30 40 50 60 70 80 90 I000 PSI

0 5 10 15 20 25 30 35 40 TONS/IN 2

0 10 20 30 40 50 60 KG/MM"YIELD STRENGIH

Fig. 5- - Fracture toughness inde:r diagram for 1-inch thick alu-minum alloy plates representative of commercial practice

"96

FRACTURE TOUGHNESS INDEX DIAGRAM FOR STEELS

The principal fracture toughness test method employedin NRL studies of high strength steels has been the drop-weight tear test (DWTT) as correlated with results obtainedin the large structural prototype element explosion-teartest (ETT). These methods and the correlation procedurehave previously been described in detail (1); From thespectrum investigation of 1-in. thick steels, a simplifiedFracture Toughness Index Diagram (FTID) which indexesthe DWTT fracture toughness characteristics of a steelin terms of ETT performance of the material has becaevolved. The FTID chart relating specifically to all1-ia.-thick steel plates evaluated to date is presentedin Fig. 54.

The FTID given in Fig. 54 has been updated to presentthe latest information. Previous studies have emphasizedthe point that the fracture toughness characteristics of ametal at a given level of YS are highly sensitive to processhistory, thickness, etc, in addition to chemical composition.A rational evaluation of process history effects is depictedin Fig. 54 by the various curves designated "optimum mater-ials trend line" (OMTL). These curves separate the dataillustrated into characteristic groups relating to millprocessing variables (melting practice and/or cross-rolling)of the steels. The OMTL may be recognize( as the reference"yardstick" for evaluation of new steels with respect to theprac'ticable upper limits of fracture toughness for any givenYb level as a function of conventional or special processingvariables. All steel data relating to the limiting ceilingOMTL curve have involved vacuum-induction-melt (VIM) or adouble vacuum practice of VIM plus vacuum-arc remelt and theillustrated ceiling OMTL curve represents the current esti-mate of the optimum limits attainable with special VIM and1 to 1 cross-rolling practices.

The significance and interpretation of the FTID datagiven in Fig, 54 will not be repeated here since they havebeen covered in detail in Refs. 2, 7, 24, 25, and 27.Wherever possible, new data (to be described) ave reportedin charts containing the basic curves and correlationfeatures of the FTID for 1-in, thick steels to illustratea ready comparison of results for the new steels withpreviously tested and reported Cata.

97

0 0c

w w Oii Nt.44_j~ 0-

w W 114,1L~wow i w 2 0 1. V4j

w- 0 I 0 Ue 0

C-0 u> > 0 W>t )t V ri - 0 I

>._ WI.f - =

LL / E-4 41.1E-4 .

z r0I >.- -- w uz (

OZ0j 0. 8 0 0LJN0 0 - .,4 %

(8--d A9.N uVIIO3 00d

z < W980000 c -, o-0 >

DROP-WEIGHT TEAR TESTS (DWTT) OF 5Ni-Cr-Mo-V STEEL PLATE[HY-130(T)]

A total of twelve 80 ton electric furnace heats (halfof which were vacuum degassed) of the 5Ni-Cr-Mo-V steelwere produced. Plates ranging from 3/4 to 4-in. thickfrom six of these heats (including the first and eleventh)considered to be representative of the production experienceof this material have been studied at NRL. The tensileand fracture toughness properties of these plates aregiven in Table 16.

A graphical summary of the data given in Table 16is depicted in Fig. 55, as referenced to the FTID featuresfor 1-in. thick steels. For the 1 and 2-in. thick HY-130(T)plates, (open symbols in Fig. 55) the standard DWTT energyvalues are noted to range from approximately 3500 to 5000ft-lb which is essentially equivalent to the DWTT valuesdeveloped by 1-in. thick HY-80 hull steels. (See Fig. 54.)In all cases the DWTT fractures for the 1 and 2-in. thicksteels were found to exhibit full shear (450 slant fracture)characteristics (Fig. 56) at 30 0 F.

The data given in Fig. 55 for the 2 to 4-in. thickplates represent DWTT values obtained for specimens ma-chined to 1-in. thickness.

The 2400 to 4200 ft-lb range of standard DWTT energyvalues obtained for the 3 and 4-in. thick HY-130(T) steelsstill implies a relatively high level of fracture tough-ness. This may be observed by noting the ETT correlationindex levels of plastic strain for fracture propagationgiven by the basic FTID chart, At this stage of develop-ment, it is not intended to emphasize the specific valuesobtained for YS and DWTT energy of these steels or thatthe variability depicted in Fig. 55 in through-thicknessproperties is necessarily representative of the expectedproperties in all "properly heat-treated" 3 and 4-in. thickplates of the HY-130(T) steel.

The results of this study emphasize the need formethods and criteria for evaluation of heavy section plates.This need was recognized, and the DWTT was devised as alaboratory tool for measuring the full thickness fracturetoughness by integrating the through-the-plate-thicknessquality gradients. Full thickness DWTT of a variety of2 to 4-in. steels are underway and equipment, instrumenta-tion and specimens are being assembled to facilitate thetesting of pressure vessel steels with thicknesses up to

99

i!

TABLE 16

TEST DATA FOR HY-130(T) STEEL ALLOY (INi-Cr-Mo-V) PLATES[All specimens in "weak" (WR) orientation]

Thick- 0.305" Diam Tension Test Data Charpy V Drop-Weight TearNRL ness 0.21 YS U El.in 2" RA 30"F * 306FNo. (in.) (kei) (ksi) (M) (%) (ft-lb) (ft-lb)

H14 2/4 161.3 158.3 18.0 61.3 83 2958

H13 : i 145.0 153.5 18.7 63.5 30 4558

398 1 )93.7 141.5 18.8 64.0 93 >5000

J21 1 134.8 140.1 18.9 64.1 79 4636

386 1 138.3 146.3 18.3 58,5 76 4478

J93 1 138.4 144.7 17.5 58.8 54 3775

J94 1 137.8 ±45.3 18.0 60.1 - 4596

395 1 137.3 144.1 18.0 60.4 - 3966

396 1 136.7 145.3 17.5 69.9 - 4052

J97 1 137.2 146.7 18.0 64.5 - >5000

398 1 137.1 145.9 18.0 64.2 - 4494

Kul 1 136.5 143.0 18.7 63.7 57 --

K20 1 137.6 143.2 19.0 64.4 54 --

K21 1 137.4 156.8 17.5 59.0 52 --

G73 2 144.9 156.7 18.0 64.6 55 3557

G74 2 149.2 157.1 18.0 59.9 71 3843

H12 2 146.7 158.3 18.3 58.3 70 4175

B99 2 141.8 148.6 17.5 60.2 66 3775

375 3 140.6 147.7 18.0 61.1 65 4188(Top Surface)143.3 149.4 12.5 37.6 55 35 8 4(Centerline)120.2 150.5 17.5 56.4 44 3609(Bot.Surface)

J74 4 151.3 155.2 17.0 55.2 43 2618(Top Surface)143.9 153.7 17.5 52.7 53 3 3 3 3(Centerline)140.7 156.6 17.0 56.2 42 2412(Bot.Surface)

*No. H14 tested as 3/4-in.-thick by standard DWTT specimens;all 2, 3, and 4-in.-thick steels were split and tested asstandard 1-in.-thick DWTT specimens.

100

HL M

uI ct - - ý- F t

LC) w L

cr~~(/ 03a.a40 a

0- Q JSli.4-

>1 0, '1-

L'4 Ua

0~ 0- 0

o~~ ~ u

(0~~~ 0C ~ N~ a4e,

101

F440~

HY 130 PLAl E 137 YS 3966 DWTT

Fig. 56 - Drop-weight- tear test fracture faces of a 1-in.thick HY-130(T) steel alloy (J-95) illustrating full shear(450 slant fracture) characteristics found to be typical

for 30OF standard DWTT for 1 to 2-in, thick H-Y-130(T)plates

12-in. Preliminary results obtained for 2 and 3-in. thickHY-130(T) steels are described below.

FULL THICKNESS DWTT OF 2-3-IN. THICK HY-130(T) STEELS

The single pendulum impact machine used at NRL formaking direct energy absorption measurements on DWTTspecimens is equipped with replaceable tups that providefor a capacity of either 5,000 or 10,000 ft-lb. The smallercapacity tup has proved to be adequate for the testing ofmost 1-in. thick steels for fracture propagation in the"weak" (WR-orientation) direction (See Fig. 54). However,the 10,000 ft-lb capacity machine has been required for"strong" (RW-orientation) direction tests of 1-in. thicksteels, and also for the 2 to 3-in. thick steels charac-terized by "low" to "moderate" fracture toughness pro-perties. Heavy section steels of "high" fracture toughnessproperties could not be evaluated with the single pendulumimpact machine.

Initial studies of high toughness, heavy sectionsteels required the use of the 1 ton (22,000 ft-lb capacity)or the 6 ton (180,000 ft-lb capacity) NRI, drop-weightmachines. These early studies also required that severalspecimens be tested at a given temperature to establishthe amount of energy required for complete fracture, i.e.,by a "bracketing" technique which was the method originallyemployed (and hence the name, DWTT) when the test was firstconceived. The size of the specimens required for DWTTof 2 to 3-in. thick steels (T x 8 x 28-in.) made thematerials requirement for a single determination by the"bracketing" technique prohibitively excessive, and itwas necessary to look to other means for conducting DWTTof heavy section steels. A solution to this problem wasachieved by adapting the existing drop-weight machineswith auxiliary equipment comprising photographic equipment tand an instrumented impact transducer to record load signalon either a cathodc-iay oscilloscope or a tape recorder.This enables the force-time data to be permanently recordedduring an impact test and provides for the indirect calcu-lation of energy absorption through the relationshipsinvolving the impulse of the force acting on the specimenand the change in linear momentum. The details of thistechnique will be described in a future report.

The preliminary results of the full thickness DWTTof the 2 and 3-in. thick HY-130(T) steels are presentedin Table 17. It is apparent that both of the HY-130(T)

103

CCY

00 m - I

4-)E

t- 00 00 C)

'4.4 - - - -

CUY

4-j -4 -4-4 ~'- C I

14 O-M*' c ii c

M- 0o- o q vL

0 - M0 -4 4 00 C

~z4 -4 -4 CD 'J 1o~ ~ 1Z~ ~ -4 N- C

E-4r

0o0C0 0

Cf1

104

steels exhibit relatively high fracture toughness charac-teristics at 30*F. However, these data also suggest thatboth of these steels are within their "transition tem-perature range" for tests at 00 F. It is also noted thatthe data given in Table 17 implies that on the basis of"nominal-net-section-test-area" the 2-in. thick HY-130(T)steel (10 sq. inof "test"area) is considerably higher at30OF and essentially equal at 00 F to the DWTT fracturetoughness properties of the 3-in. thick HY-130(T) steel(15 square inches of "test" area).

The fracture appearances of the full thickness DWTTspecimens of the 2 and 3-in. thick HY-130(T) steel testedat 30°F are shown in Fig. 57. Apparent in these frac-tures are indications of centerline segregations of non-metallic inclusions which result in longitudinal rupturesin planes parallel to the rolled surfaces of the plates.This condition is partly attributable to the melting anddeoxidation practices used for the HY-130(T) steels.Extensive studies by the producer are underway to developrefinements of melting and deoxidation practices toalleviate or eliminate the tendency to promote centerlineplanes-of-weakness for fracture propagation parallel tothe rolled plate surfaces.

DROP WEIGHT TEAR TEST STUDY OF 2OvMn-2%Ni WELD METALSDEVELOPED FOR FABRICATING HY-130(T) STEEL WELDMENTS

Weldments, 2 to 3 ft. in length, representative ofeach electrode and weld procedure under development bysubcontractors in the NAVSHIPSYSCOM HY-130 steel programwith U.S. Steel were furnished NRL to investigate theirfracture toughness characteristics. These weldments areconsidtred representative of the best 2%Mn-2'Ni fillerwire and procedures developed during the first phase ofthe program. Details of the welding processes and pro-cedures used are given in Table 18. The results obtainedin this investigation are summarized in Table 19.

Figure 58 presents a summary of the 30°F DWTT energyvalues and YS relationships obtained for the 2VMn-2%Niwelds in Table 19, as referenced to the FTID for 1-in.thick steel plates. The shaded region given in thisillustration encompasses the range of DWTT-YS valuesdeveloped with 1-in, thick DWTT specimens from the 1 to2-in. thick HY-130(T) steel plates (dcscribed Previouslyin Fig. 55). It is noted that the 2%Mn-2%Ni welds arecharacterized by generally lower DWTT fracture toughnessvalues than that exhibited by the HY-130(T) plates.

41

105

Fig. 57 ilustrating the 30 cF drop-w~eight-tear -test fracturefaces of the 2 and 3-in, thick HY 130(T) steel alloys reportedin Table 17

106

0.

N e 2 N 0 a0 02q N0 4 N in LO

L1 4 C6 4 q i '1

0..

IU

z N+ • 0 0 0

LO N

+ +

00 0l 0 0 W

+ +. >_ > > >

107 r- v-

v. a l v 4c l CD to

oz Q0 10 10 0 x V x 0CI II I -P : V : vin v & 4 N 4 54

06 CL 8 8

0 - -~ -~ ' -

I ~U N 10 10 " 10'

TABLE 19

TEST DATA FOR l-IN,-THICX HY-13U(T) (5Ni-Cr-Mo-V) WELDMENTS

AFV~l'-'r Av. Cv Av. C"NRL Location 0.2% YS UTS El.in 2" RA 0 30OF 0 30 0 F 0 80OF

No. (ksi) (ksi) (%) (%) (ft-lb) (ft-lb) (ft-lb)

J86 MIG-FlatAll Weld 146,5 155.7 18 0 61.8 1720 51 60Trans. Weld* i40.0 149 4 13.0 37.4 - _Plate (WRE 139.3 146.3 18.3 58.5 4478 76 76

J96 VIG-FlatAll Weld 144.9 154.1 1,0 64.5 1723 53/65** 53/72**Trans. Weld* 141.5 150.9 11.0 63.9 - - -Plate (WR) 136.7 145.3 17.5 69.9 4052

J93 MIG-Vert.-IlpAll 'Weld - - - - 3400 62 68Trane Weld* 140 7 147 7 14.0 51.5 - - -Platp (WR) 138.4 144.7 17.5 58.8 3775 54 55

J97 MIG-Vert.-UpAll Weld 140.4 151.? 18.0 64 6 22,25 51/59** 55/66**Trans. Weld 139 2 151.9 14.0 64.3 - - -

Plate (,.R) 137.2 146.7 18.0 64.5 >5000

J98 MIG-Vert.-UpAll Weld 140.1 152.6 IS U 64 4 4000 64/83** 64/86**Trans. Weld 139 2 150 4 15.0 64 3 - - -

Plate (WR) 137 1 145.9 18.0 64.2 4494

J94 Stick-FlatAll Weld 136.5 137.9 19.0 63.6 3333 59 69Trans. Weld 134.1 138.3 15.5 60 3 - - -Plate (WR) 137.8 145.3 18.0 60.1 4596

Jc'5 Stick-Vert.-UpAll Weld 134.7 140 1 17.0 58.8 3201 54 GOITans. Weld 133.2 1,') - 16.0 63.3 - -

Plate (WR) 137.3 14.l 18.0 60.A 3966

Kl1 Stick-FlatAll Weld 140.7 144.2 17.5 60,09 1905 55 55Trans. Weld - - - - . -

Plate (WR) 136.5 143.0 18.7 63.7 - 57 57

K20 Stick-Vert.-UpAll Weld 140.7 151.6 18.0 59.0 1326 45 50Trans. Weld .- - - - -

Plate (WR) 137.6 143,2 19.0 64.4 54 54

K21 Stick-Vert.-DownAll Weld 140.2 140.2 16.0 55.8 - 53 60Tralhs. Weld - -. - - - -

Plate (WR) 137.4 156.8 17.5 59.0 - 52 52

• Tensile fracture outside of weld in Plate-HAZ area.:* Charpy V values for TOP/BOTTOM of weld deposit.

108

108

wi w

U) 0

I- Q. 1 0w D~ D

os UiCjj .0

0 ~iw~ ~J.- Ucur

5f

M F

U - -.i 0

NC z zc

4. L

'v 4

'41

w 84

o~~~r 0 0 0 04

0 0 0 0 0 0 L0

0 0 0, 0 -4o

1090

Depending upon the electrode and welding procedure, theYS values of the 2%Mn-2%Ni welds range from 134.7 to 146.5ksi; and the DWTT values range from 1326 to 4000 ft-lb.It should also be noted that a DWTT value is not reportedfor the manual-STICK weldment fabricated in the vertical-down position. Difficulties encountered in electron-beam weld preparation of specimens from this weldmentresulted in subsequent DWTT fractures involving significantamounts of the HY-130(T) plate metal, with the concomitantly"high" DWTT energy values considered to be more represent-ative of the plate rather than the weld.

The data given in Table 19 and Fig. 58 indicate anextremely large variation in DWTT values for the 2%Mn-2%Niwelds despite the fact that the chemical compositions ofthese welds were reported to be essentially similar. Ofparticular significance is the observation that the MIG-spray-arc flat-position welds are characterized by rela-tively low (approximately 1700 ft-lb) DWTT values at 30'Fcompared with the significantly higher DWTT values (approxi-mately 3400 ft-lb) obtained with the same heat of weld wireused in MIG-interrupted arc-vertical up position welds.On the basis of Cv results for these 2%Mn-2%Ni welds(described below) the "low" DWTT values for the MIG-Flatwelds cited above, and the similarly "low" DWTT resultsobtained for some of the manual-STICK welds depicted inTable 19 and Fig. 57 are considered to reflect transitiontemperature range aspects of these weld metals.

CHARPY V (Cv) FRACTURE TOUGHNESS OF HY-130(T) PLATES AND2%Kn-2%Ni WELD METALS

The average C energy absorption curves for 1 and2-in. thick HY-130TT) steels are depicted in Fig. 59.It should be noted that these data represent tests in the"weak" (WR orientation) fracture direction. The 1-in.thick steels exhibit Cv shelf level values ranging from52 to 93 ft-lb at 30 0 F. In general, the shelf level valuesare maintained for these steels to test temperatures atOF and for several 1-in. thick steels to subzero testtemperatures. The one 2-in. thick steel which does notexhibit Cv shelf level energy values at 30OF (No. G-73,Fig. 59, bottom) represents data for material cut fromone end of a 2-in. thick plate produced from the firstproduction heat of HY-130(T) steel. In the productionof this plate, a malfunction of the heat treatment andquenching facilities was encountered and reported by theproducer; but the specific plate was shipped without re-quiring the plate to be recycled through the continuous

110

!00 14"-I" HY 130 (T) STEEL........... H98

80 _ ,H14H 1380_K21 -_--

-J8

L 60 0

r 0 HY 130 (T) STEEL

G 73

>-60, 7

40

40-

20 30 0F

-160 -120 -80 -40 0 40 80

TE MPERATURE (0F)

Fig. 59 - Summary of the average Charpy V ICy)energy absorptioni

curves for 1 and 2-in, thick H-Y-130(T) steel plate alloys. The data

relate to wweakw (WR orientation) fractulre direction tests.I

LdG7

heat treating line to t nsure the development of optimumproperties. Laboratory heat treatment of small platesamples of the G73 plate resulted in the development ofCv properties essentially equal to those shown in Fig.59, bottom, for the G74 plate which also represents another2-in. thick plate produced from the first production heatof HY-130(T) steel.

The Cv data developed by NRL for the 3 and 4-in.thick HY-130(T) steel plates are given in Figs. 60 and61. It is noted that the significant variation in surfaceand centerline properties previously described for thesesteels by results of standard DWTT are similarly indicatedby the illustrated results for Cv tests. Both the 3 and4-in. thick steels are characterized by Cv shelf levelvalues at 30 0 F. However, depending on the location andorientation of the Cv specimen, it is apparent that Cvshelf level values below the 50 ft-lb minimum shelf levelenergy specified as the aim for HY-130(T) steels aredeveloped by the "improperly" heat treated 3 and 4-in.thick HY-130(T) steels.

The average Cv energy absorption curves for the various2%Mn-2%Ni MIG and STICK welds studied in this investigationare depicted in Figs. 62 and 6ý. All weldments were fabri-cated with a wide gap (1-in.), single V joint preparation(required for DWTT of weld metals) as shown in Fig. 62,top, right. This results in the evaluation of essentiallyundiluted weld metal. As noted in the illustration, Cvtest specimens were machined so as to study top and bottomportions of the 1-in. thick weldments. For those weldmentsfabricated with back-up strips (Nos. J96, J97, J98, K21),dilution of the weld metal with the HY-130(T) steel isnoted to result in relatively higher Cv energy values forthe "bottom" specimens compared with the Cv values obtainedfor undiluted weld metal in the top portion of the welds.

Generally, all of the illustrated 2%Mn-2%Ni weldmetals are noted to develop average Cv values of 40 ft-lbor more in tests at 30*F. However, the characteristicshape of the Cv energy absorption curves indicates thatseveral of these weld metals are within their transitiontemperature range at 30 0 F; and the relatively low DWTTvalues described previously for these weld metals wouldseem to confirm this fact. This raises a question con-cerning the O°F fracture toughness characteristics ofthese weld metals if the HY-130(T) weldment is used ascontemplated in surface ship applications to gain desiredfabrication experience prior to use in submarines.

112

-4

3-IN.-THICK 5Ni-Cr-Mo-V HY 130 STEEL PLATE (J75)

80- 080T8--TOP SURFACE

RW-ORIENTATION gA CEERLN60 . xBOTTOM

w 20-3"

(D -30"F

z 0-

0

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•I X2 BOTTOMS-" • -• ""SURFACE-40 - -

x

30OF

o fI I I I I I I I I I I 1-120 -80 -40 0 40 80

TEMPERATURE (OF)

Fig. 60 - Charpy V energy absorption curves for both surfaces andcenterline locations in a 3-in. thick HY-130(T) alloy (No. J-75)

113 1

4-iN-THICK 5Ni-Cr-.Mo-V HY 130 STEEL PLATE (J74)

80-

RW -ORIENTATIONACENTERLINE

60 -- -... -- -BOTTOM

0x SURFACE

40 ",- TOP SURFACE

- 2F- 2- •4"

(D 306F-

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T R CENTER LINE

... _..o ,•-,•-,,=•SURFACE

40 F 6 G y n.r curves TOP SURFACEbt

20 --.

-120 -80 -40 0 40 BO 120TEMPERATURE (a F)

Fig. 61 -Charpy V energy absorption curves for both

surfaces and centerline locations in a 4-in. thick HY-130(T) alloy (No. J-94)

114

80J97-MIG -INTERRUPTED ARC .-

VERTICAL UP "X 0 J96

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60 - x

00

-80 -40 0 40 80 -90 -40 0 40 80 120TEMPERATURE (*F)

Fig. 62 -Charpy V energy absorption curves for 2%Mn-Z%Niweld metals deposited by MIG and STICK processes in HY-130

(T) steel plate

115

- ____-

K I I-MANUAL- STICK60 FLAT

x0

x

40

- 20-20 HT IP0419 36 F

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040-

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VERTICAL DOWN x

Qj 60- 0

u40"

20-HT IP0419 30OF

- 140.2 YS I01

-80 -40 0 40 80 120 160 200TEMPERATURE (OF)

Fig. 63 - Charpy V energy absorption curves for 2%Mn-2%Niweld metals deposited by STICK processes in'HY-130(T)steel plate

116

All of the HY-130(T) weldments described herein werefabricated with filler metals from large (7-1/2 to 17 ton)production heats of vacuum melted and air melted 2%Mn-2%NiMIG wires, and covered (STICK) electrodes were producedon a production basis. Initial quantities of the "production-lot" electrodes were found to be lacking in one or moreaimed for "optimum" characteristics, e.g., useability charac-teristics, root pass cracking difficulties, small deviationsfrom aimed for composition, etc. These problems weregenerally ameliorated by reprocessing and vacuum annealingof the filler metals to achieve better surface conditionson the MIG wires (free of oxide and/or die-lubricant con-tamination) and by modifications of the coating or fluxformulations used on the STICK electrodes. The generalexperience gained with these initial production heats of2%Mn-2%Ni weld metals is expected to result in the pro-duction of "improved" versions of the 2%Mn-2%Ni fillermetals which will be studied extensively along with therecently available "matching" 5%Ni-Cr-Mo-V filler metal.

DWTT OF "NEW" 9Ni-4Co-0.20C ALLOY HIGH STRENGTH STEELS

Fracture toughness data for several grades of therecently developed Ni-Co family of alloy compositions(designated 9-4-0.XXC) were presented in previous investi-gations (22, 23, 24). These have involved plates of thetwo basic compositions (available commercially sin-e 1962)containing 0.25%C and 0.45%C as well as several "experi-mental" versions of the basic alloy containing lower carbonand increased chromium and molybdenum contents from thatspecified for the basic 9-4 alloys (AMS 6540 throughAMS 6546). The proprietary designation of 9-4 may bepartly misleading in that the steels of this alloy familyare all reported to contain 7.0 to 8.5%Ni and producedto a specification range of 7.0 to 9.0%Ni. The basic9-4 composition can be adjusted or the heat treatmentvaried to develop YS levels ranging from approximately150 to 250 ksi as required for specific applications. $

Composition modifications or "rebalancing" of the alloyby the producer to secure better combinations of strengthand toughness than, that developed by the commerciallyavailable grades of the 9-4 alloys resulted in a newmember of this Ni-Co alloy family, designated 9-4-0.20Cand aimed at 180 ksi minimum YS, being made commerciallyavailable in the latter half of 1966.

In all cases, commercial production of the 9-4steels have involved special melt practices involvingvacuum arc remelt (VAR) practice of electrode ingots

117

produced from a basic electric furnace (EF) air melt heat.The initial EF melting is performed without the use ofstrong deoxidizer; and the resulting electrode ingot productis relatively high in oxygen content in the form of oxides.In subsequent vacuum arc remelting at sufficiently lowpressures, the carbon acts effectively as a deoxidizercapable of reducing oxygen to very low levels providingthe oxygen is not tied up as highly stable oxides. Theproducer has designated the term "VAR-C-DOX" to signify"vacuum-arc-remelt, carbon-deoxidation." During vacuumarc remelting, a carbon "boil" is experienced and resultsin the evolution of CO gas during the remelting operation.The specific special melt practices used for the 9-4steels are aimed at producing material essentially "free"of deoxidation product inclusions ("dirt") and of a con-siderably lower gas (02, H2 and N) content than that pro-duced by conventional practices. Present VAR facilitiesare limited to electrode ingot sizes from approximately5 to 15 tons. Usually ten (10) or more electrode ingotsare poured from the initial EF (70-80 ton) melt; however,the producer considers the product from each VAR ingotas separate "production" heats of the 9-4 alloy.

Small nlate sections (approximately 1 x 18 x 36-in.)from two of the VAR heats of material produced from thefirst large (70 ton) EF heat of the "new" 9-4-0.20C alloysteel were received in the mill quenched-and-tempered(Q&T) heat treatment condition for evaluation studiesby NRL. It is apparent from the Cv curves establishedfor the "weak" (WR) and "strong" (RW) orientations offracture in these plates, Fig. 64, that both plates hadbeen produced with relatively high (essentially 1 to 1)degrees of cross rolling. The average 0.505-in. diametertension test data also shown for these steels in Fig. 6Windicate that both steels developed YS levels slightl)abovec thz 180 ksi minimum YS specified for these steels.Fracture appearance ratings of thye steels could not beestablished from macro features of the Cv test specimens;however, the energy absorption curves suggest that Cvshelf energy values are developed by these steels at tem-peratures of OF and highei. The procurement of additionalplates of "new" 9-4-0.20C steels is planned to rievelopthe DWTT energy absorption versus temperature behaviorof these steels and, thus, establish the significanceof the relatively high (in excess of 50 ft-lb) Cv valuesnoted in Fig. 64 for subzero temceratures.

A summary of the DWTT-YS relationships for all ofthe 9-4 steels evaluated to date is presented in Fig. 65.

118

I-IN. THICK 9-4-20 STEEL PLATESYS UTS E'_ . Aý

80- No. K-28 181.3 194.5 16.2 61.

mx x

10 - x -0-j -8-o 0_4- -

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Also shown in this illustration is a comparison of the chem-ical compositions specified for the basic 9-4-0.25C (AMS 6546)alloy with that of the "new" 9-4-0.20C steel. Both of thesegrades of the 9-4 alloy are designed to develop 180 ksiminimum YS. Except for carbon content, the compositions arenoted to be similar; howeve-, significantly higher contentsof Ni, Co, Cr, and Mo are specified for the "new" 9-4-0.20Calloy than that expected for "average" composition of the9-4-0.25C basic alloy. The specific mill heat analysesgiven in the illustration for the two plates (Nos. K28 andK29) evaluated in this investigation are noted to representlow carbon, low residual Mn, Si, P, S, and high Ni, Co0 Cr,Mo, V versions of the "new" 9-4-0.20C alloy. The similarityin composition of these plates from two VAR "production"heats suggests that the basic composition of the initialEF melt predominates and controls the resulting compositionof the various VAR ingot heats produced from a given EF melt.

The "new" 9-4-0.20C steels (Nos. K28 and K29) are notedin Fig. 64 to be characterized by considerably higher DWTTfracture toughness than any of the previously tested 9-4-0steels for yield strengths of approximately 180 ksi. Pre-viously studied experimental versions of this "new" alloy(designated 9-4-0.19C and 9-4-0.23C) contained considerablylower Ni (6.81 and 7.13% respectively) and higher Cr and Mocontents than those now illustrated for the "rebalanced"alloy. The "experimental" lower nickel and carbon c'ontentand higher Cr-Mo content versions of the basic 9-4-0.25Calloy are noted to exhibit lower YS levels and higher frac-ture toughness than those obtained by the basic 9-4-0.25Calloys. At a 180 ksi minimum YS level, the DWTT fracturetoughness values of various basic 9-4-0.25 steel platesranged from 1000 to 2200 ft-lb enmpared to approx4-"1-ly3300-330C f-,b foi L. ,w i-4-n.2ue alloy stee..(Nos. K28 and K29). The high carbon content (9-4-0.45C)alloy steel is noted to be capable of propagating fracturesat elastic stress levels for all heat treatment conditionsstudied except that of Q&T 1100 0F. For this heat treatmentcondition, however, the YS of the 9-4-0.45C alloy is notedto be below the 180 ksi minimum YS level developed by thebasic 9-4-0.25C and the "new" 9-4-0.20C steels.

DWTT OF 9Ni-4Co AND MARAGING STEEL WELD METALS

The 9-4-0.25C steels and the 12-5-3 maraging steelshave appeared promising for nominally 180 ksi minimum YSstructural applications. Weld wires for these alloys, how-ever, have not been generally available commercially. Pre-liminary evaluations of producer-fabricated, TIG process,180 ksi YS weldments using "experimental" electrodes produced

121

from small (150 lb) research quality heats were reported tobe highly encouraging for weldments in both of these alloysteels (24, 25). For further studied required for the charac-terization of critical properties of such weldments, weldwires representative of semiproduction heats (1000 lb) wereprocured. Specifications for these weld metals were basedupon producers recommendations to obtain the "best" weldwire product possible. The 9Ni-4Co filler metal procuredfor this investigation was melted to an aim chemical compo-sition representing that of the "best" of the experimentalfiller metals developed for welding the basic 9-4-0.25C alloysteel as described recently by Ries and Poole (29). Thecomposition is similar but lower in Ni and Co to that ofthe new composition specified for the 9-4-0.20C alloy steel.

As described in earlier reports (5, 7, 24), two weldmetals have been evolved for fabricating the 12Ni-5Cr-3Momaraging steels. These comprise a nominal 12Ni-3Cr-3Mo(12-3-3) wire composition, expected to develop a 160 ksiminimum YS upon aging, and a nominal 17Ni-2Co-3Mo (17-2-3)wire composition expected to develop a 180 ksi minimum YSupon aging. Preliminary studies with manual TIG weldmentsindicated that both of these weld compositions developedoptimum Cv toughness-strength relationships upon aging forrelatively long times (10 to 24 hr) at 850OF and that shorttime aging at 900*F resulted in the development of YS levels15 to 20 ksi below the 160 and 180 ksi YS "expected" respec-tively for the 12-3-3 and the 17-2-3 weld metals (7). Forcomparison with continuing weldment evaluation studies ofthese two weld metals, a limited amount of a "matching" com-position filler metal (containing l2Ni-5Cr-3Mo-0.3Ti-0.3AI)was procured from the INCO Research Laboratory facility.This filler metal was produced by converting sections of2-in. thick rolled plates from a VIM heat of the 12-5-3 alloyinto weld wire.

A series of enlarged (1-in. wide) single "V" joint pre-paration weldments were fabricated by NRL to provide materialfor DWTT of the various filler metals described above. Theautomatic tungsten-arc-inert-gas (TIG) cold wire feed processwas used for all welding in this study. Generally, producers'recommendations for optimum welding procedures were followed;and these resulted in similar conditions and weld heat inputs(44 to 46 kilojoules per inch) for both the 12-5-3 maragingand the 9-4-0.25C Q&T steels. Argon at a flow rate of 40cfh was used as shielding gas (no auxiliary shielding wasused) and all weldments were started without preheat andfinished with 200-225OF maximum interpass temperature control.

122

The maraging steel weldments were aged after fabrication.A summary of the data developed in this investigation is givenin Table 20. Figure 65 presents a summary of the DWTT-YSrelationships for these high strength weld metals, asreferenced to the FTID-OMTL chart for 1-in. thick steelplates.

From the summary of data given in Table 20 and Fig. 66,it is generally noted that the 12-3-3 weld metal is charac-terized by the lowest strength and highest DWTT fracturetoughness values. The 9Ni-4Co weld metals are observed todevelop significantly higher YS levels and concomitantlylower DWTT fracture toughness levels than the 12-3-3 andthe 17-2-3 maraging steel weld metals. Surprisingly, the"matching" composition 12-5-3 weld metal which displays50 ft-lb Cv energy at 30°F is noted to develop the lowestDWTT energy value for all of the weldments evaluated.Additional C specimens of the 12-5-3 weld metal were testedat 1200 and 1 60°F with resulting values of 63 and 68 ft-lurespectively. The implications for Cv data and low DWTTvalue for the 12-5.-3 matching composition weld metal suggeststhat this weld may be within its transition temperature rangeat 30 0 F. Examination of the Cv data given in Table 20 forall maraging steel weld metals suggests that the long timeaging treatment at 850°F may be responsible for developingtransition temperature range aspects in the maraging weldmetals.

The data developed in this investigation for the 17-2-3,.Ji 'ietal, Table ?n tends to oppose th,' !caclablin drawnpreviously that optimum strength-toughness relationshipsare developed by the 17-2-3 weld metal by long timc agingat 850 0 F. It is noted that 3 hr aging at 900°F resultedin the development of YS levels slightly lower than the 180ksi YS "expected" for the 17-2-3 weld composition with DWTTtoughness values essentially equivalent to those developedby long time 850OF aging treatments. Continuing weldmentevaluation studies in this program are expected to explorethe effects of long time 8500 and 900OF aging treatmentson transition temperature behavior of the maraging weldmetals and to establish weld metal properties of maragingand 9Ni-4Co weldments fabricated with lower weld heat inputsthan those used in this investigation.

123

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RFFERENCES

1. Pellini, W. S., et al., "Review of Conccpt!. and StUtu.3of Procedures for Fracture-Safe Design of ComplexWelded Structures Involving Metals of Low to Ultra-High Strength Levels," NRL Report 63_3. June 1965

2. Goode, R J., et al., "Metallu:cgical Characteristicsof High Strength Structural Materials (Ninth QuarterlyReport)," NRL Report 6405, November 1965

3. Goode, R. J. and Huber, R.W., "Fracture ToughnessCharacteristics of Some Titanium Alloys for Deep-DivingVehicles,"Journal of Metals, August 1965

4. Hc-, T)C, , "Ffffacts of Heat Treatment on the Stress-Corrosion-Cracking Resistance of Several Ti-Al-Mo-VAlloys", Report of NRL Progress, February 1966, pp,14-18.

5. Goode, R. J., et al., "Metallurgical Characteristics ofHigh Strength Structural Materials (Tenth QuarterlyReport),"NRL Report 6454, April 1966

6. 'Howe, D. G., "Effect of Heat Treatment on the Stress-Corrosion-Cracking Resistance of Several Titanium Alloys",Report of NRL Progress, July 1966, pp. 20-21

7. Puzak, P.P., et al. , "Metall uigical Characteristics ofHigh Strength Structural Materials (Eleventh QuarterlyRepour"•,NRL Report 6513, August 1966

8. Howe, D. G., °'Properties of Ultra High Strength Alloys;(Effect of Heat Treatment on the Stress-Corros-;ion-C-cackingResistance of Several Titanium Alloys)", Report of NRLProgress, September 19(6, pp. 25-27

9. Howe, D. G., "Eiffects of Heat Treatiag Environmental Con-ditions on the Stress-Co,-rosion.Cracking Resistance ofSeveral. Titanium Alloys". Report ol N3UL Progress,February 1967

10. Brown, B. .. and Reachem, C. D., "A Study of the SiressFactor in Corrosion Crackinrig by Use of the PrecrackedCantilever Beam Specimen", Corrosion Science, V1. 5,No. 11, 1965, pp. 745-730

12E6

11. Kies, J. A., et al., "Fracture Testing of Weldments" inFracture Toughness Testing and Its Applications, ASTMSTP 38 , American Sc- ety for Testing ,-an Materials,1965

12. Brown, B.F., "A New Stress-Corrosion Cracking TcstProcedure for High Strength Alloys," ASTM MaterialsResearch and Standards, Vol. 66, No. • M-ari-c-h7-9-6--

13. Brown, B. F., et al., "Marine Corrosion Studies, (ThirdInterim Report of Progress)" NRL Memo Report 1634, July1965

14. ASTM Materials Research and Standards, Vol. 4, No. 5,Ma:_y 19 6

15. Judy, R. W., Jr., and Crooker, T. W., et al., "Fracto-graphic Analysis of Ti-7A1-2Cb--lTa and Ti-6A1-4VFractures Developed in 'Wet' Fatigue," NRL Report 6330,January 1966

16. Lange, E.A., and Klier, E.P., "A Study of FractureDevelopment and Materials Properties in PVRC Vessels1 and 2," Welding Journal Research Supplement, Vol. 27,No. 2, February 1962

17. Lange, E.A., et al., "Failure Analysis of PVRC VesselNo. 5," Welding Research Council Bulietin, No. 98,August 1964

18. Pellini, W.S., and Puzak, P.P., :Fracture Analysis Dia-gram; Pi~cedAures for the Fracture-Safe EngineeringDesign of Steel Structures," NRL Report 5920, March 1963

19. Kooistra, LF. , Lange, E.A. . and Piukett, A.G., "Full-Size Pressure Vessel Testing and Its Application toDesign," "TRANS of the ASMF Journal of Engineering forPower, VoI.- SRei, s A-No. t -

20. Crook-or, T.W. , and Lange, E. A., "Low Cycle FatigueCrack Propagation Resistance of Materials for LargeWelded Structures", Proceedings., ASTM Fatigue CrackPropagation Sympo,;ium'••--AYt-li-ii- City. N.J., June 30-July 1, 1966 (publication pending)

21. Crooker, T.W., i'nd Lange, EA. , "Corr.s1,mn FatigueCrack Propagation in Modern High PFe-formo..n,.o Struc Lura,_Steels." AIME Sympesimw, The Effect o)f Enviro(nnent onthe Mechanicai Properties of ',!etais, Chicago. Illinois,October 31, 1966 (publication pending)

127

22. Goode, R.J., et al., "Metallurgical Characteristics ofHigh Strength Structural Materials (Fourth QuarterlyReport)", NRL Report 6137, June 1964

23. Puzak, P.P., Lloyd, K.B., Goode, R.J., Huber, R.W.,Howe, D.G., Ciooker, T.W., Lange, E.A., Pellini, W.S.,"Metallurgical Characteristics of High Strength Struc-tural Materials, (Third Quarterly Report),"NRL Report6086, January 1964

24. Puzak, P.P., Lloyd, K.B., Goode, R.J., Huber, R.W.,Howe, D.G., Judy, R.W., Jr., Crooker, T. W., Morey, R.E.,Lange, E.A., and Freed, C.N. , "Metallurgical Character-istics of High Strength Structural Materials, (Eigh'nQuarterly Report)", NRL. Report 6364, August 1965

25. Goodc., R.J., et al., "Metallurgical Characteristics ofHigh Strength Structural Materials, (Seventh QuarterlyReport)", NRL Report 6327, May 1965

26. The ASTIM Committee on Fracture Testing of High-StrengthMetallic Materials, "The Slocv: Growth and Rapid Propaga-tion of Cracks"(Second Report), ASTM Materials ResearchSzndarAa• Vol. 1, No. 5, May 1961

27. Pellini, W.S., et al., 'Metallurgical Characteristicsof High Strength Structural Materials, (Sixth QuarterlyReport), N"RL Report 6258, December 1964

28. Pellini, W.S. and Puzak, P.P. "Factor-, that Determinethe Applicability of High Strength Quenched and TemperedSteels to Submarine Hull Construction", NRL Report 5892,December 5, 1962

29. Ries, G.D. and Poole, S.W .. W'elding Journal ,c-ýearchSupplement. October 1966, Vol. 31, No. 10, p. 465-S

128

Si, ,nt) C is~,fs, atnn

DOCUMENT CONTROL DATA R & D

Naval Research Laboratory UnclassifiedWashington, D. C. 203901G- 0-,1

I R1 9ORT t

METALLURGICAL CHARACTERISTICS OF HIGH STRENGTH STRUCTURALMATERIALS (Twelfth Progress Report)

4 OEICPI TO4F NOTESF (7'Yp,0Oe torIP IOfet inCl daltes)

Progress report on the project5 ýu T O 1OR,$I (First nonti* middl. ,nltI, 1.8f

Goode, R.J., Huber, R.W.. Howe, D.G., Judy, R.W.,Jr., Crooker, T.W.,Lange, E.A., Freed, C.N., and Puzak, P.P.RE-OP T DATE ?A1. TOTAL N1O OF PAGES lb 'O O RIE'S

September 1967 133 29go C aN TR•IACT 01 O4 GRAiN T NO M OP.GI'lA REPORT V• RLIPOR

NRL Problem Nos. FOI-17SRoJEý, T. M01-05 NRL Report 6607

MOI-18

M 03-01 960 To r;.rR RIPOt NOj1i (AO: -th., -6.b..1 ai d,.ri,ii be *...ld.d

M04-08Bd.

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.Special Projects-Department of the NavyOffice of Naval Research

[Advance Re.,earch Projects Agencc

A progress report covering the research studies in high strongth structu-ialmetals c )nducted during the period July 1966 ti.rouph __ 2inaary 1967 is presente..The report includes fracture toughness stuci z-.5 on sonie new titanium alloys as *w--1ias results of an initial feasibility study air- ud at raising the optimum strength andtoughness limits for titanium alloys through --n approach involving thick sectioncom-posites. The results of salt water streq-s-corrooion-cracking stL.dies on a numberof titanium alloys are described includ ng a study concerning the effect o" vacuum-nheat treatment on stress-corrosion-ci tc`Kirng resistance of tUtanium alloys. Fatigý.lecrack propagation studies in air arid in salt water on the pressure vessel steelsA30IB, AZ01B, and A517F are discussed; and the results cot-.pared to the actualservice performance of these samne ti .tturrals in PVRC program. std(i~es. Prelin'-.: 1-

ary fracture toughness correlation diay.,anms are prcserted for .igh strength st - ibased upon fracture mechanics and erg-necring test methods. The latest versioro,the Fracture Toughness index Diagrams for steels, titanium alloys and Aluminu•malloys, a usual feature of this report series and which are based upor engineeringtest methods, are presented and the significance of the features ri-fiy disc.saedFracture toughness studies on thick plates of 5Ni-Cr-Mo-V, and d "new" ý`Ni-4,Ccsteel are reported. In addition, the results are presented for similar stulies onwelds of 2Mn-2Ni in 5Ni-Cr-Mo-V steels and welds of 9Ni-4Co, 12'Ni-3Cr-3.Mc an17Ni-ZCo-3Mo compositions.

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Security Classification

14 LI. NK A LINK I U LIN8 CKEY KOROS - -

MOLE WT ROLj *1 ROLIE T

Higi -strength structural metalsFracture toughness studiesHigh..strength titanium alloysHe?'. treatmentsStress-corrosion-cracking resistance of

titanium alloysLov cycle fatigu'e crack propagation in

pressure vessel steelsAluminum alloys

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