cobalt monograph series - IAEA

cobalt monograph series


cobalt-base superalloys - 1970 november 1970

cobalt alloy permanent magnets december 1971

cobalt-containing high-strength steels September 1974

Cover

FrontReplica electron micrograph of HP 9-4-n steel after martensiticquenching, showing self-tempered martensiie with cementite parti-cles.

BackThin-foil electron micrograph of 13 Ni (400) maraging steel aged for4 hours at 900" F (480° C), showing martensite laths; their granularaspect is due to & very fine precipitate, probably a-FoMo.


cobalt-containing high-strength steels

a critical review of the physical metallurgy of cobalt-containing high-strength steels, and a survey of theirprocessing, properties and uses

A. MAGNEE J.M. DRAPIER J . DUMONT D. GOUTSOURADIS L. HABRAKEN

Centre de Reoherches MStallurglques, Centre d'Information Centre de Recherches Mitallucgiques,Li&ge, Belgium du Cobalt, Brussels Liege, Belgium

CENTRE D'INFORMATION DU COBALT, BRUSSELS

1974

Gentre d'information du cobalt s.a.

Centre d'lnformation du CobaltRue Royale 66 B-1000 Bruxelles (Belgique).

Cobalt Information Center, c/o Battelle Memorial InstituteKing Avenue 505 Columbus, Ohio 43201 (U.S.A.).

Cobalt Information Centre7 Rolls Buildings, Fetter Lane London EC4A IJA (England).

Kobalt-InformationElisabethstrasse 14 D-4 Diisseldorf (Deutsdiland).

FOREWORD

The first commercial cobalt-containing high-strength steels of the carbide-hardened, maragingand stainless types date back to 1960-1961. They provide a classic example of an industrial break-through based partly on intuitive reasoning, in lhat the final products turned out to exhibitproperties which surprised the industrial world, and perhaps even the scientists responsible fortheir development. Although the introduction of these steels was obviously preceded by severalyears of laboratory research, the origin of their outstanding properties remained unexplainedfor quite a time. For example, in his conclusions to the Journees Internationales des Applicationsdu Cobalt, held in Brussels in 1964, the undersigned, reflecting the general consensus of bothspeakers and audience, had to admit that the role of cobalt in these steels, and more particularlyits remarkable strengthening potential when associated with molybdenum, was still obscure.

Since then, a considecable amount of information has been generated on cobalt-containing high-strength steels, and further grades have been developed. It progressively became apparent thatthe difficulty encountered when trying to explain the role of cobalt lay in the fact that its effect,though major, is indirect. In particular, the contribution of cobalt to solid-solution hardeningis small, and it is not involved directly in the formation of strengthening precipitates. However,it has a favourable effect on the martensite formation temperature and refines the martensiticstructure; it also exerts a decisive influence on the precipitation kinetics, favouring the formationand retention of fine precipitates whose presence results in considerable strengthening; finally,it can participate in an ordering process. Needless to say, much of the new insight into thephysical metallurgy aspects of these steels was made possible by the availability of examinationtechniques of greatly increased sensitivity.

The importance of these findings, as well as the growing industrial significance of cobalt-contain-ing high-strength steels, prompted the Cobalt Information Centre to devote the third volume inits " Cobalt Monograph Series " to presenting a critical summary of the knowledge on the threegroups of steels under consideration. The task of writing the manuscript was entrusted to theCentre de Recherches Metallurgiques; as was the case for the preceding volumes in the series,the physical metallurgy aspects of these materials were emphasized, but processing, propertiesand uses were also dealt with, though more concisely.

This monograph is based, not only on a comprehensive literature survey, but also on theexperience acquired at C.R.M. during its long-standing association with C.I.C., as well as onnumerous discussions with specialists in this area of metallurgy. The authors gratefully acknow-ledge the active support received from the following scientists who, in addition to participatingin the discussions, also provided initial readings of the manuscript: Dr. J.H. Gross andDr. S.J. Matas (Chapters III and TV on carbide-strengthened steels); Dr. S. Floreen (Chapters Vto VII on Ni-Co-Mo maraging steels); Dr. H. Brandis, Dr. R.L. Caton, Dr. A. Kasak, Dr. A. vonden Steinen and Dr. D. Webster (Chapters VIII and IX on stainless maraging steels). They alsowish to thank their colleagues both at the Centre de Recherches Metallurgiques and the Centred'Information du Cobalt for xheir help during the preparation of this volume. In particular,they wish to acknowledge the invaluable assistance of Mrs. H. Lefebvre, of the Brussels C.I.C.office, who assumed a largo part of the editorial work, from manuscript finalization to proofchecking.

The authors have attempted to present a coherent but concise review of the vast amount ofinformation available to date on the three families of alloy steels. They believe that this volumewill prove useful to scientists concerned with the strengthening mechanisms of high-strengthsteels, to metallurgists involved in promoting them, and to engineers in their continuing searchfor new materials capable of meeting the increasingly stringent service conditions imposed bypresent-day technology. If this monograph stimulates a reader response either to identify thoseaspects which require further clarification or to suggest ways of improving properties, thenprogress will be assured and the authors' task adequately rewarded.

Professor L. Habraken

contents

iNlRHIHfllOV. '

SlKINi.IlitNISi, MttllAMSUS IS HkiH-S IKfcN'li IH SfEf.LS . . 2

2 i . Solid-Solution Strengthening -

2 I. Phase-Transformation Sirensthening : Bainilic Reactions . 42.2.\. Morphology of Bainites 42.^.2. Properties o( Bainiies ft2.2.3. Role of Cobalt 7

2,.;. Phu>e-Transformation Strengthening : Martensiiic Reactions 72 .".I. Cieneral Charaaenstics of Martensiiic

"(ran>t'ornii:iions 7I . ' . - . Types of Mancnsile 10: . : O . Morphology of Lath Martensitc 10I..v4. Morpholoyy of Twinned Manensite II_ . ? 5 . Transition from Lath to Twinned Martensite. . . . 122.}.b. Properties of Martensites 132.3.7. Controlled Martensitic Transformation 15

2.-I. Precipitation Strengthening 162.4.!. Mechanisms 162.4.2. Carbide Precipitation 18

2.4.3. Precipitation of Intermetallic Compounds . . . . 19

2.5. Strengthening b> Thermomechanicat Treatment . . . . 20

C \RBii>t-STRt\t;iHfcNi:i) STF.EI.S — PHYSICAL METALLURGY . . 22

3.1. Introduction 22

}.2. Continuous Cooling and Isothermal Transformations . . • 23

3.2.1. Continuous Cooling Transformation (CCT) Curves 233.2.2. Isothermal Transformation Curves 24

3.?. Biiinitic Transformation Structures 243.3.1. Bainites Formed on Continuous Cooling . . . . 243.3.2. Mechanical Properties Associated with

Continuous-Cooling Bainites 273.3.3. Isothermal Bainite in HF 9-4-45 273.3.4. Mechanical Properties of Isothermal Bainite in

HP 9-4-45 27

3.4. Martensitic Transformation on Quenching 28

3.5. Tempering Reactions 293.5.1. HP 9-4-X Steels 293.5.2. 5Ni-Cr-Mo Steels 313.5.3. lONi-Co-Cr-Mo Steels 323.5.4. Retained Austenite 34

3.6. Effect of Alloying Elements on Tempering Response,Strength and Toughness 343.6.1. Effect of Carbon 253.6.2. Effect of Nickel 353.6.3. Effect of Silicon and Manganese 363.6.4. Effect of Carbide-Forming Elements 363.6.5. Effect of Cobalt 373.6.6. Strength/Toughness vs. Structure Relationship . . 38

3.7. Concluding Remarks 39

4. CARBIDE-STRENGTHENED STEELS — PROCESSING AND PROPERTIES 40

4.1. Primary Processing 40

4.2. Properties , 424.2.1. Strength/Toughness Characteristics 424.2.2. High- and Low-Temperature Properties 464.2.3. Fatigue Behaviour 474.2.4. Stress-Corrosion Characteristics 47

4.3. Secondary Processing 48

4.4. Applications 49

5. Ni-Co-Mo MARAGING STEELS — PHYSICAL METALLURGY . . 50

5.1. Background 505.1.1. Role of Alloying Elements 525.1.2. Compositions 53

5.2. Martensitic Transformation 545.2.1. Formation and Morphology of Martensite . . . . 545.2.2. Factors Controlling Lath Martensite Formation. . 56

5.3. Ageing of Martensite 575.3.1. Precipitation Reactions 57 .5.3.2. Ordering 605.3.3. The Cobalt/Molybdenum Interaction 605.3.4. Maraging Kinetics 62

5.4. Austenite Reversion 64

5.5. Strength/Toughness vs. Structure Relationship 67

6. Ni-Co-Mo MARAGWG STEELS — THE CONVENTIONAL GRADES 68


6.2. Properties 686.2.1. Strength/Toughness Characteristics 686.2.2. High- and Low-Temperature Properties 726.2.3. Fatigue Behaviour . ' , . . . ' 746.2.4. Stress-Corrosion Characteristics 74



7. Ni-Co-Mo MARAGING STEELS — THE ULTRA-HIGH STRENGTH

GRADES 77

7.1. Processing 77

7.2. Properties 797.2.1. Strength and Toughness 797.2.2. High- and Low-Temperature Properties 797.2.3. Other Properties 80

7.3. Applications .80

8. STAINLESS MARAGING STEELS — PHYSICAL METALLURGY . . 81

8.1. Background 818.2. Effect of Alloying Elements on Equilibrium Structures . . 838.3; Transformation Temperatures and Structures 85

8.3.1. Martensitic Transformation 858.3.2. Austenite Reversion 888.3.3. Retained Austenite 89

8.4. Grain Size 90

5.5. Ageing Reactions 91ii.5.1. Iron-Chromium System 91S.5.2. Fe-Cr-Co and Fe-Cr-Ni Systems 928.5.-1. Fe-Cr-Co-Mo and More Complex Ailoys . . . . . 948.5.4. Concluding Remarks 98

8.6. Snengih Toughness vs. Structure Relationship 995.6.1. General ' 9 95.6.2. Effect of Retained Ausienite 998.6.3. Effect of Delta-Ferrite 1008.6.4. Effect of Prior Austenite Grain Size 101

9. STAINLESS MARAC.ING STEELS — PROCESSING AND PROPERTIES 101


9.2. Properties 1039.2.1. Strength/Toughness Relationship 1039.2.2. High- and Low-Temperature Properties and Thermal

Stability 1079.2.3. Fatigue Behaviour 1099.2.4. Corrosion Resistance 1109.2.5. Stress-Corrosion Characteristics . . . . . . . . I l lV.2.6. High-Tempemiure Oxidation Resistance 113



10. CONCLUSIONS 114

REFERENCES 116

AUTHOR INDEX 124

SUBJECT INDEX 127

I. INTRODUCTION

The importance of high-strength steels in modern technology is demonstrated by thetremendous effort, which has been devoted to the development and understanding o'i thisclass of materials over the past twenty years. In Figure I, the yield strengths of steelscurrently used in so-called '• massive structural applications ", on the one hnnr1 and in" specialized structural applications", on-the other, are plotted versus decade from1850 onwards. This graph illustrates both the progress that has been mac!,- over the vcarswith respect to strength of structural steels and the impact exerted by the advent of highlysophisticated constructions such as aircraft or missiles on current ;ind future strengthdemand, it shows, in particular, a sharp increase in the slope of the curve lor thespecialized applications as from about 1950, and the anticipated coming into use of steelswith strengths of 400,000 psi (2700 MN/m-) or higher by 1980.

The introduction of cobalt-containing steels exhibiting very high strength a! roor,,temperature dates hack a mere fifteen years or so, but these steels have achieved almonimmediate notoriousness because of their outstanding properties. The purpose of th smonograph is three-fold : (1) to present as complete a survey as possible of existingand developmental cobalt-containing high-strength steels; (2) to attempt to explain dierole of cobalt in these steels through an exhaustive examination of their physical met?Iluigy:(3) to situate, whenever possible, these steels within the overall group of present-dayhigh-strength steels. As in our previous monographs, the emphasis will be put or. therelationship between properties and structure, although more practical data are alsoprovided. Cobalt-containing high-speed steels and heat-resistin? steels will not be deal!with in this volume, since their fields of application have little in common with that ofhigh-strength steels.

The subject matter is divided into four main sections. Chapter 2, which is a general reviewof the hardening mechanisms operative in high-strength steels, is intended to provide therequired background for the subsequent sections.

Chapters 3 and 4. are devoted to carbide-strengthened steeis, essentially those of the9Ni-4Co and !0Ni-8Co-Cr-Mo types. Their physicalmetallurgy is dealt with in Chapter 3.while Chapter 4 provides a summary of their processing, properties and uses.

In Chapters 5, 6 and 7, the .Ni-Co-Mo maraging steels are described, first from thephysical-metallurgy viewpoint (Chapter 5), and then from a practical aspect. On accountof conspicuous differences in the type and amount of information available on the" conventional " and " ultra-high-strength " grades, it was found convenient to treatthese two classes in separate chapters (respectively Chapters 6 and 7).

Fig. 1.1.—- Increase in yield strengthof steels used in structural applica*tions during the ymrs since iS50.

After A.M. 1-ALL [/./].

400

350

J.30OSx-2506§200

ft 150

g 100

SO

~~l TComponents for engines,,air frames, missiles, .and other specializedstructural application;.,

Lreomotives.bridges,' i >^ships, buildings, TVtowers, —jr—and other massive L ^structural applications

1900YEAR

1950

1

I OB-U !-t liMMNIMi HIGH-SIRE NCiTK STF.EI.S

l-'in;.ll>, Chapter.-. v and V are devoted to maraging stainless steels, all of which containcobalt and molybdenum, in addition to chromium for stainlessness. Here again, the formerchapter deals with their physical metallurgy, and the latter with thsir properties and uses.

The references quoted in the text are listed at the end of the book. In order to keep this\olumc to •>. reasonable size, their number was voluntarily restricted. Selection of thereferences retained vwis based on criteria of conclusiveness and originality, but the reader>hould appreciate thai this practice has resulted in the omission of numerous papersof \alue in confirming the findings of other workers. Where no specific reference is cited,it should be assumed that the information presented is derived from the experience gainelat the Centre de Recherche* Metallurgiques during its long-standing involvement in workon steels in general and on cobalt-containing ones in particular.

:. STRENGTHENING MECHANISMS IN HIGH-STRENGTK STEELS

The aim of this chapter is to re\iew briefly the transformation and precipitation mechanismsthat are operatise in high-strength steels, and to discuss their influence on the latter'smechanical properties. As in earlier review papers [2.1 to 2.3], the emphasis will be placedon the effect of cobalt on these mechanisms, reference being made in each case to thesroups of steels more particularly concerned. The various classes of steels discussed inthis monograph are listed in Table 2.1. which also shows the strengthening mechanismsinvolved for each group and the chapters in which they are described.

2.1. Solid-Solution Strengthening

If an alloying element in steel does not form alloy carbides or intermelallic precipitates,then it will be in substitutional solid solution in the iron lattice, which may be either ferriteor austenite. As such, it will exert a solid-solution hardening effect which may be usefulin increasing the strength of the steel. Theories on solid-solution strengthening [2.4]assume that the solute atoms are not distributed uniformly but form " segregations '*.The latter are actually responsible for the solid-solution hardening, since they raise theenergy required to move dislocations.

Segregation of atoms can occur through interaction between solute atoms and imper-fections, or between the solute atoms themselves. Elastic interaction between soluteatoms and dislocations leads to Cottrell atmospheres, whereas interaction of electricalorigin between these atoms and stacking faults gives rise to the Suzuki effect. Finally,

TABLE 2.1. — STRENGTHENING MECHANISMS INVOLVEDIN HIGH-STRENGTH STEELS UNDER DISCUSSION

Steels

Low-alloy steels

Cai bide-strengthenedmartensi'-ic steels

(9Ni-4Co. 10Ni-8Co-Cr-Mo)

Ni-Co-Mo maraging steels

Stainless maraging steels

Strengthening Mechanisms

Martensitic transformation (twinned martensite)Precipitation of carbides

Bainitic or inartensitic (lath or twinned) transformationPrecipitation of carbides

Martensitic transformation (low-carbon lath martensite)Precipitation of intermetallic compounds

Martensitic transformation (essentially lath martensite)Precipitation of carbides and/or intermetallic compounds

Chapter

2

3,4

5,6,7

8,9

2. STRENGTHENING MECHANISMS IN HIGH-STRENGTH STEELS

ATOMIC % SOLUTE ATOMIC V. SOUJTE2 4. 6

ATOMIC'/.SOLUTE

Fig. 2.1. — Effect of elements in solid solution on the R.T. lattice parameter,modulus of elasticity and shear modulus of a-iron. After W.C. LESLIE [2.5].

interaction between the solute atoms themselves results in short-range ordering (Fisher)or clustering. As will be seen in Chapter 5 (Section 5.3.2), the occurrence of an orderingreaction in a cobalt-containing solid solution probably contributes to the strengtheningof maraging steels on ageing. Some neutron-diffraction data have indicated that hardeningof these steels may be due to a low degree of long-range order or a high degree ofshon-range order.

Data on both lattice dilatation and the changes in elastic constants with solute concen-tration are required in order to estimate the magnitude of the solid-solution strengtheningeffect. The changes in the room-temperature lattice parameter, modulus of elasticity andshear modulus of a-Fe brought about by the presence of various solid-solution elementsare illustrated in Figure 2.1. It is seen that additipn of up to about 6 at. % Co has anegligible effect on the lattice parameter, whereas both the modulus of elasticity and theshear modulus increase with increasing cobalt additions. Figure 2,2, which is a plot ofthe size misfit parameter, za = (l/a0) (da/dc), versus the change in shear stress with soluteconcentration, AT/Ac, shows that little strengthening of a-Fe occurs for low concentrationsof cobalt. This confirms earlier experimental work which had shown that the solid-solution hardening effect of cobalt in ferritic and austenitic steels is, respectively,approximately 5 and 4 HV per I wt.% Co [2.6].

Fig. 2.2. — Correlation of R.T. solid-solution strengthening of iron-basealloys with size:, misfit parameter.

After;^V.G;-;LESLIEJ[2J]/: .

- 5"g

i*|9NI.

Be.

Ptr-30

002 0.01 006 ODB D.1D 0-12 OK 0.16 0

i H l l i H S I R l N C I I H S i r U S

The strengthening effect of elements in substitutional solutions can be influenced, some-times markedly, by the interstitial content of the solid solution, as well as by more readilycontrollable variables such as substructure, grain size and temperature [-.7], Work on aO.lC-l2Cr martensitic steel [-.*'] has shown that, for cobalt additions up to approximately15",,. a hardness incrjmeni of approximately 9 HV per 1 vvt.% Co occurs. The highersolid-solution hardening effect of cobalt in u murtensitic structure can be used to advantagein high-strength mi.rtensitic steels, as will be seen in Chapter 3 (Section 3.6.6).

2.2. Phase-Transformation Strengthening : Bainitic Reactions

The simplest way of achieving high strength in a steel is to make use of the lower trans-fcimaiion temperature structures, buinite and martensite. In this and the followingsection, only the more general characteristics of such transformations will be Hisciissed.As regards bainitic reactions, detailed discussions will be ,'ound in recent publications(.\y to".".//].

2.2.1. Morphology of Bainites

The following definition of bainite will be used in this Monograph : bainite is a constituentof steels which is formed by the decomposition of austenite within a temperature rangelocated between the field of martensite formation and that of ferrite and peariite formation.This constituent consists of an aggregate of ferrite and carbides or partly stabilizedaustenite. Its morphology changes progressively with the transformation temperature, in

C-enriched zones

Nucleation ;ind growthof upper bainite

Transformationfranl

Austenite-marlrnsite nodule -•• Acicular ferrile'

Nudcation and growthof lower bainite

Formation of massive bainite showing gran,uiar (left) and acicular (right) aspects.

a) C >0.3% ( v 10,000) b) C < 0 . 2 % (X 2200)

Fig. 2.3. — Morphology of bainitic structure in high-carbon steels (>Q.3%)and Sower-carbon steels (<0.2%, generally with alloying elements).

2. STRENGTHENING MECHANISMS IN HIGH-STRENGTH STEF.I.S

that the size of the particles increases with decreasing temperature, as does the acicularityof the structure.

As illustrated in Figure 2.3, three distinct types of bainite, designated massive, upper, andlower bainites, can be produced in many steels by the athermal transformation of austeniteat suitable cooling rates or by isothermal transformation at subcriiical temperatures;the first is also designated granular baiciite (when observed at low magnification) oracicular fcrrite (when observed at high magnification). In the isothermal transformationdiagrams for steels with sufficiently high carbon contents, the bainitic field is roughlydivided into two horizontal bands corresponding to upper and lower bainite [2.9]. Forupper bainite, which forms above 660-750°F (35O-4OO°C) depending on the steel's com-position, ferrite plates nucleate within the austenite grain; carbon diffuses av ay andconcentrates in the remaining austenite until it finally precipitates as Fe3C carbide betweenthe ferrite plates (Fig. 2.4a). The carbides are often parallel to the axis of the ferriteneedles and the cementite-ferrite orientation relationships depend on both the cementite-austenite and austenite-ferrite orientations.

For lower bainite, i.e., that obtained at lower temperatures, the ferrite piates nucleate inthe same way, but carbon no longer diffuses so readily and cementite, occasionally precededby e carbides, precipitate within the ferrite plates (Fig. 2.4i.'). The carbides form anangle of approximately 60° with the axis of the ferrite needles. Moreover, it has beenshown that the cementite-ferrite orientation relationships in lower baiuite are identicalwith those that prevail in tempered martensite; examples pertaining to 9Ni-4Co steelswill be found in Chapter 3 (Section 3.3.1).

Massive bainite forms easily in low-alloy steels during athermal transformation ofaustenite; it corresponds to microstructures consisting of coarse plates or presenting analmost entirely granular aspect (Fig. 2.5). It occurs over a range of cooling rates that isdependent on the steel's composition, and consists of irregular ferrite grains having amoderately high dislocation density and islands of enriched austenite which may trans-form either completely or partly io martensite [2.9]. In order to explain both thestabilization of the austenite and the characteristic appearance of massive bainite, it has

t

a) Fully transformed at 930oF (500°C) :upper bainite.

b) Fully transformed at.750°F (400°C) :lower bainite.

Fig. 2.4. — Microstructural features of isothermal bainitic transformationin low-carbon steel. After K.J. IRVINE [2.12]. <-i-•••":' ''• x 16,000

Fig 2 i —' Microstructural1 features ofathermal bainitic transformation in 0.1C-lCr-0.5Mo-0.002B sleel fully transformedat 885°F (475°C). After M.E. BUSH andP.M. KBLLY [2.13]. x 1500

O > H U I i O N i U M N O U K . H - S I R l N l . l t l N f l - l U S

been sueuested that dehomogenization of the austenite occurs on holding the latterin us meuistahle temperature range [2.14], Carbon would diffuse under ihe effect of anacmitv gradient, as opposed to a concentration gradien1. towards regions having a highden>i\> of imperfections. As a result, piior to its transformaiion. the austenite wouldcontain a network of carbon-rich regions. From this stage onwards, the theory for theformation of upper bainite can be applied, to a first approximation, to massive bainite(hiah temperature of formation, low cooling rate) : nucleation of supersaturated ferriteplatelets within a carbon-depleted region, followed by their growth and the possiblerejection of the carbon, towards tht untransformed austenite. The fairly large size of thehainuic ferme areas in massive bainite results from the more rapid growth of the plateletsdue to the prior dehomogenization of the austenite. The austenite islands are eitherswallowed up during growth or imprisoned between two ferritic regions. Their stabilizationcan be attributed to their high carbon content, which is due partly to the initial inhomo-senett> of the austenitc. and panK to the possible subsequent diffusion process; thedislocation densit) can also be a stabilizing factor.

2.2.2. Properties of Bainites

Both the isothermal and continuous-cooling (atherma!) b.iinites have been investigatedextensively over the past few years, in relation to the development of ultra-high strength,low-alloy steels. Alloys that can be air cooled to bainitic structures have acquired greatercommercial importance and have consequently formed the subject of m;;ny of the recentinvestigations concerned with the microstructure / mechanical property relationship inbainitic steels.

The various factors which are thought to contribute to the strength of upper- and lower-bainite type materials produced by air cooling have been described qualitatively [2.10.2.11]. It has also been postulated [2.9] that similar factors should apply to massive bainites.More recently [2.IJ], it was shown that the latter owe their base strength to contributionsfrom the strength of pure iron i.nd the effects of alloying elements in solid solution, theP'ior austenite grain size, iiod the dislocation density. The strength is raiser1 above thebase level by the presence of the islands referred to in the preceding section, through atwo-phase effect similar to that which leads to the strengthening of fetnte-pearlite aggregatesby pearlite colonies. Isothermal bainites §we their base strength to the same factors asmassive bainites but with an increased dislocation-density contribution, the acicularferrite matrix being rich in defects. The strength is raised above the base level by furthercontributions from the dislocation density, which increase as the transformationtemperature decreases. The carbide precipitates also exert a large strengthening effectwhich is dependent on both the number and shape of the carbides present. As regardstoughness, it is known that, for numerous steels, lower bainites are tougher than upperbainites.

Since the microstructure of the steels is controlled by the transformation temperature, arelationship exists between strength and transformation temperature. This is illustratedin Figure 2.6. The diagram shown in Figure 2.7, on the other hand, indicates the com-bination of yield strength and impact transition temperature which can be obtained inferritic-pearlitic, bainitic, and quenched-and-tempered low-carbon steels. The micro-structure of the latter consists of acicular fernte together with small precipitated carbideswhich are more finely dispersed than in the low-temperature bainitic structure [2.12].

It is also possible to illustrate in a single diagram (Fig. 2.8) the important strengtheningfactors in these three types of microstructure. The base line in this diagram represents arelationship between yield strength and grain size. For the C-Mn-Nb steels, strengtheningbeyond that related to grain size is due to carbide precipitation. In the case of bainiticsteels, the strength increase above the yield strength / grain size line depends upon both

2. STRENGTHENING MECHANISMS JN HIGH-STRENGTH STEELS

Temperature" ol maximum rale of (rarefontHlton, t„ . too son _6QO TOO eoa

VELD. STREN5TH,'MN/m*

• MAR- .

TENSIIE!BAINITE M FERRITE]

1 IPEARUTEI200u-iooo..noo woTemperature dt maximum rate of lransformailon,.*F

Fig. 2.6. — Relationship betweentransformation temperature andstrength of low-carbon bainiticsteels. After K.J. IRVINE [2.12\.

••~KT.Tr:.:..- ~~ ./•"- 8 0 ;YIELD STRENGTH; »> psi.

Fig. 2.7. — Relationship between yieldstrength and impact transition tempera-ture of low-carbon steels for differentmicrostructures. After KJ. IRVINE [2.12].

.5 . 10 15 20GRAIN SIZE. d''4 mm' ' 4

Fig. 2.8. — Relationship between grainsize and yield strength of low-carbonsteels for different microsiructures.

After K.J. IRVINE [2.12].

the dispersion of the fine carbide particles and the dislocation density in the bainitic ferriie.Finally, as regards martensitic low-carbon steels, their structure is not greatly differentfrom that of lower bainites, since their high martensitic start temperature causes self-tempering to occur. However, the carbide particles are finer than in the bainitic structure,the dislocation density within the ferrite plates is high, and some carbon remains in solutionin the martensite; hence, the block representing the martensitic structures lies just abovethe highest strength level for the bainitic structures.

2.2.3. Role of Cobalt

Cobalt additions are not normally made to bainitic steels, since this element scarcelyinfluences their hardenability and grain size, and does not exert any appreciable solid-solution hardening effect on ferrite. At most, when present in large amounts togetherwith other alloying elements, cobalt might promote strengthening by leading to a betterdispersion of the carbides, or to the formation of intermetallic compounds. Also, in somelow-alloy high-strength steels, addition of small amounts of cobalt may decrease thebrittleness [2.5]. However, there is an important group of cobalt-containing steels inwhich the bainitic transformation is used to advantage, viz. the 9Ni-4Co steels, which arecharacterized by high strength and exceptional notch toughness (cf. Chapters 3 and 4).

2.3. Phase-Transformation Strengthening : Martensitic Reactions

As was the case for the bainitic reactions, the present discussion will be restricted to areview of the general features of the martensitic reactions. A fuller treatment of the subjectwill be found in References 2.15 to 2.19.

2.3.1. General Characteristics of Martensitic Transformations

When steel is cooled rapidly from the austenitic region, transformation occurs at lowtemperatures and this leads to the production of a new phase, martensite. As opposed toa conventional nucleation-and-growth reaction, the martensitic transformation ischaracterized by the fact that it cannot be suppressed by rapid cooling, even using veryhigh cooling rates (up to 32,500°F/s, i.e. 18,000°C/s). In addition, it requires no diffusion

. O H A L I ' - l O M M M M i H K i H S l K l \ ( . I H S T t l l S

or tiuermixinsi of the atoms involved in the phase change; the transformation productsinherit ihe composition of the parent phase and each atom lends to retain iis originalneighbours. Finally, martensiiie reactions are displaeive or shear-like, in that the atomsmo\e co-operatively to produce substantial shape changes in the transforming region,e\en thouah the movement of each atom is small compared with the atomic diameter.A consequence of this dilTusionless transformation is ihe magnitude of irk* strain associatedwith the transformation from the parent to the product phase, which reaches as muchas 10",, in ma,lv commercial alloys. An experimental indication of these transformationstrains i> given by the upheavals that occur on a polished surface following transformation.

The followinc laws are obeyed by most martensitic transformations [-.-W] :(i) On cooline. the transformation begins at a temperature A/» which may vary as a resultof the prior thermal and mechanical history of the material but is generally iiniepcmlcnini tin- («<>/i"ji!j rate. For highly alloyed ferrous materials, the martensitic transformationtray also be induced at temperatures above the A/« point by plastic deformation; thehighest temperature at which this can occur is known as the M,i temperature.

lii) When the nucleation event is insensitive to time, the transformation proceeds primarilywhile the temperature drops below .W», giving rise to the well-known atlwrmal type ofmartensitic reaction. However, there are instances in which ivjcleation may occur isu-ihermalh: and the martensitic plaits will then exhibit rapid growth even at a fixedtemperature.

(iii) Both the >ize of the martensiie grains and the A/.< temperature are dependent on theausienite grain si/e, which itself is a function of the time and temperature of the austenizingtreatment : the smaller are the austenite grains, the smaller will be the martensite grains.

(iv) The extent of the transformation can be decreased by using a very high cooling rateor holding the alloy within a specific temperature range. Stabilization of austenite incarbon steels is dependent on this characteristic.

(v) A temperature hysteresis exists between the heating and cooling reactions; thetemperature at which marlensite starts to transform to austenite on heating at the rateof 4.5 F min (2.5 C mint is known as the A* point.

Of the models put forward to explain the nucleation and growth of martensite, the mostacceptable to date is that which proposes that embryos of the manensitic phase formwithin the austenite. and that some of these reach a critical size which enables them togrow at temperatures beiow Ms [2.21]: the small regions of martensite thus formed areseparated from the austenite phase by a semi-coherent interface. A recent structuralinvestigation of the 18Ni(3OO) maraging steel [2.22] did in fact show that, on cooling belowA/«. martensite platelets nucleated within the austenite grains themselves and notpreferentially at the grain boundaries; on the other hand, their growth was stopped eitherby the prior grain boundaries or by previously formed platelets. Moreover, the plateletsdid not appear to nucleate and grow uniformly within the structure, since some regionswere almost completely transformed to martensite, while others were still entirelyaustenitic.

There is considerable interest in the effect of alloying elements on the M, and As

temperatures. The A/., temperature is important because it can affect the resultingmechanical properties of the steel. In low-carbon, low-alloy steels, the M, temperaturelies between 570 and 750 F (300 ;ind 400 C) and full transformation can be obtained oncooling through the martensite transformation range to room temperature. Whenappreciable amounts of alloying elements are present,-however, the Ms temperature isdepressed to near room temperature and the M, temperature, i.e., that at which themartensitic transformation is complete, may be well below room temperature. In this case,the steel will contain retained austenite which can affect the maximum strength obtainable.

8

: STRKNGTHRNINC; MEl"NANISMS IN HIGH-STKLNGTH STF.ELS

l-iy. 2.9. -- Influence of cobalt on the .1,anil XI, temperatures of various steels.

Fe-22.5Ni After R.B.G. YEO [2.24\K--|7Ni-l.5Mo ( After G.W. TLTFNF.LI.Fc-|9Ni-l.5Mo| and R.I.. CAIRNS [2.25]Fe-l2Cr-4Ni After CM. HAMMOND [2.26\l-c-1 KCr-O. IC After I"). Cot JTSC UKADIS [2.27]Fe-13Cr-O.7C After D. CtiLTSOLRAOis

and I.. HAIIKAKHN [2.2S]

5 10 15Co CONTENT wt.V.

The influence of cobalt on the A/,, temperature of steels may vary greatly. In lo\v-all.,\steels, a coefficient of -• 22 F ( - 12 C ) per weight percent Co has been reported [2.P3].Data on more complex compositions are shown in Figure 2.9. It is seen that additionof up to 8l,,Co to an Fe-22.5Ni alloy raises the M* temperature. This effeci applies toFe-Ni alloys and maraging steels in general; it may result from the lowering of theaustenite shear modulus by the cobalt addition [2.29]. However, a direct relationshipbetween Ms and cobalt content is not always observed. For instance, the Xis temperatureof Fe-18Cr and Fe-12Cr-4Ni steels decreases with increasing cobalt contents. Figure 2.9also shows that increasing the cobalt content of Fe - 17 to 19Ni - L5Mo alloys from 15 to19.5% lowers their A/* temperature; this is in contrast to the behaviour observed in Fe-Nialloys and maraging steels for cobalt contents below 8%. Data on the effect of otherelements on the Ms temperature of steels are given in Chapters 5 and 8 (Sections 5.2.2and 8.3.1).

The correlation between cobalt content and Ms temperature can be used to control themicrostructure of an alloy steel. As the ailoy composition becomes more complex, thesteel may become borderline with respect to both the possibility of avoiding ferrite in themicrostructure and of keeping the martensite transformation range close io roomtemperature. If a conventional austenite-forming element such as nickel is used tocounteract the ferrite-forming tendency, then the martensitic transformation range may bedepressed below room temperature. On the contrary, if cobalt is used, the structure canbe controlled without depressing the A/s temperature.

As regards the As temperature, it is strongly depressed by addition of appreciablequantities of alloying elements such as nickel; this has a direct bearing on the choice otthe tempering temperature. Cobalt also has a depressive effect, though much weakerthan that of nickel1 this effect is retained in chromium steels (Fig. 2.9), but in Fe-Ni alloysand maraging steels in general, cobalt has been shown to raise the />„ temperature.

t O l U l | - u » \ l VIM M i Hl l . i l s ( R I A ( , 1 H S I M I S

\

\

Fig. ; 10 • Microsiructural features if lalh manensiie in Fe-1.94Mo alloy. X 14,000After G. KRALSS and A.R. MARDER [2.17).

2.3.2. Types of Martensite

Two major types of martensite form in iron-base alloys as a result of the shear-type,ditTusionless rnariensitic transformation of austenite. The gradual recognition of thesetwo martensitic products, which differ with respect to composition, range of formationin a given alloy system, crystallography, morphology, and fine structure, has led to theuse of a multiplicity of terms to differentiate between them. A definite preference forthe term " lath martensite " for one form and " plate martensite " for the other wasrecently reported [2.17]. Although the former will be used in this monograph, an alternatedesignation, viz. " twinned martensite ", has been preferred for the latter.

2.3.3. Morphology of Lath Martensite

Figure 2.10 is a typical example of the structure of lath martensite as observed in low-carbon steels. The optical and electron micrographs show that the basic units, whichare planar and lie along one direction, are generally aligned parallel to one another ingroups that have been termed packets, fragments, blocks or sheaves. These packets arethe predominant feature of the microstructure and the individual martensitic laths arevisible as a fine substructure within the packets. Several packets are found within a prioraustenite grain. The widths of the units which make up a packet of martensite range fromless than one tenth to several microns, with the most frequently occurring width beingbetween 0.1 and 0.2 \xm. Adjacent martensite laths may be separated by low- or high-angleboundaries or may be twin-related [2.30, 2.31], but this does not appear to be generallytrue [2.32].

The orientation relationship between martensite packets and the austenite parent phaseare generally of the Kurdjumov-Sachs type, i.e. :

(ii i). ,//(oil)* Uio]T / / r jTi] ,The interface between martensite packets and the austenite is parallel to a j l l l ) r typeplane [2.32]. Moreover, the orientation relationships between adjacent laths formingthe same martensite packet are generally of the type :

(1 ">)»,//(110)«2 [TllW/[001]a2

Lath martensite is also characterized by its substructure which consists predominantlyof a high density of tangled dislocations within the laths. It has been stated thatdislocations in structures of this kind tend to lie in <11 l>a directions and are predominantlyscrew in nature [2.33].

10

2 STRI- MECHANISMS IN HIGH-STRENGTH STEELS

The lath martensitic transformation differs from the twinned manensitic transformationin carbon steels mainly in that the martensite packets grow more slowly and the changesin the austenite-martensiie interface with time are not the same. The latter involve avery short-range diffusion process, over two 10 three atomic layers, in order to relievethe stresses se: up by ihe large volume change thai occurs during the transformation[2.22], Lath martensite formation in binary iron-nickel alloys containing around 10 to25 "„ Ni is of particular importance within the framework of this monograph since it islargely responsible for the properties of the maraging steels to be discussed in Chapters 5lo 7. The high density and quasi-uniform distribution of dislocation.* in this structurefavour a belter age-hardening response (see Section 2.4) by providing a large number ofnuclcation sites and increasing the diffusion rales, hence ensuring a more uniformdistribution of finer precipitates.

2.3.4. Morphology of Twinned Martensite

Twinned martensite differs from lath martensitc in that adjacent plates do not formparallel to one another. The first, plates formed lend to span their parent austenite grainsand effectively partition the austenite, thus limiting the size of subsequent plates, as shownin Figure 2.1 la. The effect of this partitioning is to produce a wide range of plate sizesin this type of martensite. A characteristic midrib is evident in most of the plates.

The substructure of twinned martensite consists partly of fine parallel twins (Fig. 2.1 \h).In most cases, the mode of twinning is found to be j 112]M type [234\, although recentlythe 11 IOjM type was observed for an Fe-1.82C alloy [2.35]. Frequently a fine dislocationstructure consisting of parallel arrays of screw dislocations is found to coexist with finetransformation twins in plates of Fe-Ni martensite [2.33].

In this monograph, the twinned martensitic transformation is of interest in the discussionof carbide-strengthened martensitic (Chapter 3) and stainless maraging (Chapter 8) steels.

a) in Fe-1.86C alloy, showing auto-catalylically formed plates. * 440

b) in Fe-32Ni alloy, showing twinnedsubstructure of plates. x 40,000

Fig. 2.11. — Microstructural features of twinned marlensiie. After G. KRAUSS and A.R. MARDER [2.17].

11

H U I i i : \ l \ l \ l \ i . I 1 1 1 . I I S I U ! N u l l ! S U M S

2 . . \ 5 . Transition iron: !.alh (<> Twiniit'J Marti >t*iic

Several explanations for the transition from lulh to twinned nuiriensite have beenproposed The occurrence of one or the mher type is controlled. ;il least partly, by thecarbon content of the -steel: as shown in r is;ure 2.12. the martensue of plain-carbon steelscontaining up 10 U.4",,C consists mainly of lath nutrtensile. but the ratio of lath totwinned nianenOie decreases rapidl> for higher rarhon conteills. On the other hand,the morphological charnic which develops in Ke-C alloys with increasing carbon contenth.is been aitrihiiied to twin formation below a critical A/, range of 570 to 425 1- tU)0 to22II O \2 .U]. Another stud} has led to the conclusion that not oniy temperature butal-o allo\ composition may signiticantly intluence the extent of twinning in the marlensiteI - V i : the important factor appears to he the relative magnitudes of the critical resolvedshear -tresses (or twinning and slip at a giv.'n temperature for a particular alloycompx-iiion

It h.is ,i!>o been suggested that low slacking-fault energy (SI lit favours the formation oflath nuirtensite [2Jx]. However, there are addition elements which have oppositeeffects on the SH of austeuile. yet favour the formation of twinned martensite when addedin increasing amounts [-./"]: this >- ''he ease for manganese, which lowers the SFE. andnickel, which raises it. According to a fourth proposal [-J-J. lath martensiie in ironalloys is alway- associated with a cubic phase, whereas twinned martensite is tetragonal.The cubic phase in l-"e-C and Fe-N alloys is assumed to be due to Zener disordering of theimerstitiul atoms as the martensite forms. The major exception to this hypothesis is lhatcubic twinned manensiie forms in Fe-Ni. Fe-Pt and Fe-Mn alloys.

In work on the structure and mechanical properties of Fe-Ni-Co-C and Fe-Cr-Co-Csteels [2.3'J and ?.4<>. resp.]. it was stated that the V, temperature is not a sufficientindication of the extent of internal twinning in martensite when comparing two differentsteels. Howe\er. for a given steel, the temperature range over which martensiie formsdetermines the extent of twinning in the martensite units. It was suggested that the drivingforce for transformation, as given by the area of the transformation hysteresis loop, maygive a better indication of the occurrence of twinning in martensite. A larger hysteresisgap would mean a greater driving force and consequently a greater chance for the martensiteunits to be twinned [2J'J. 2.41}.

From a compilation of existing data on the morphology of martensite in binary iron alloys,and consideration of the effect of alloying elements on the y field in the correspondingphase diagrams, it was also suggested that the y-stabilizcd systems form lath martensiteand. with sufficient additions, twinned or h.c.p. (e) martensite as well, whereas the y-loopsystems can transform to laih martensiie only [2.17]. This proposal lends support tothe temperature hypothesis mentioned at the beginning of this section. The limitedaustenite solid solution in the y-loop systems does not permit sufficient alloy additionsKi reduce the V/« temperature of the austenite to the critical temperature. On the otherhand, sufficient alloy additions can be made in the case of the y-stabilized systems topromote twinned martensite formation.

The effect of cobalt on iron alloys does not contradict the y-stabilization and hence thetemperature hypotheses [2.17]. In fact the y phase in the Fe-Co system is stabilized;howevet. only kith marlensilc forms because coi-ilr additions raise the A-/, temperatureso thai the critical temperature to induce formation of twinned martensite is never reached.

The work on the structure and mechanical properties of Fe-Ni-Co-C and Fe-Cr-Co-Csteels [2.39, 2.40] showed that cobalt is not effective in retaining lath martensite forincreasing carbon contents. Although the addition of cobalt raises the MH temperatureof such steels, it does not reduce twinning unless the Mx temperature is such lhat thecritical resolved shear stress for slip is lower than that for twinning.

12


RELATIVE VOLUME %LATH MARTENSITE

I

wt.%C

Fig. 2.12. — Effect of carbon content on relative volume frac-tions of lath and twinned martensites in Fe-C alloys. TheAA, temperatures and volume percentages of retained austeniteare also shown. After G.R. SPEICH and W.C. LESLIE [2.36].

ICONICKEL, at.%

2 < 6 8 10 _ B 20 25 30 35

RANDOM _ i LATH ! TWINNED

Fig. 2.13. — Separation of solid-sol-ution hardening and transformationsubstructure effects in Fe-Ni alloys.

After G.R. SPEICH and P.R. SWANN [2.31].

2.3.6. Properties of Martensites

As pointed out in [2.42], the potential effects of diffusionless transformations on strengtharise from : (i) ultrafine subdivision of the product-phase grain into twins; (ii> significantincrease in the dislocation density; (iii) re^.nement of the parent grain size in the absenceof twinned products; (iv) additional solid-solution hardening; (v) additional hardeningthrough ordering of the product phase. However, martensites are not necessarily hard andstrong, which means that the mere occurrence of a martensitic transformation does notper se confer conspicuous strength upon the product of the reaction. Even in the case ofsteel, the martensitic transformation does not result in high strength if the carbon contentis low.

In the absence of interstitial solutes, the greatest hardness that can be attained by amartensitic transformation in unalloyed iron or dilute iron-base substitutional solidsolutions is about 260 HV, with a yield strength of about 100,000 psi (700 MN/m-). Thisstrength is little affected by significant changes in the structure, such as the transitionfrom lath to t inned martensite [2.31, 2.32]. In Figure 2.13 the yield strength ofquenched and of .-ecrystallized Fe-;Ni alloys is plotted against the square root of the nickelcontent; the flow stress c -of the quenched alloys can be separated into three components :

af = (<*a + <?Ni) + GG + Gs (-')where ao is the frictional stress for pure iron, <7Ni is the increase in frictional stress causedby the nickel addition (solid-solution hardening effect of nickel), cG is the grain sizecontribution (equal to kd - ' >2 where k is the Fetch slope and d the grain diameter), and os

is the increase in flow stress due to the defect structure introduced by the transformation(this term includes the effects of the increased dislocation density, the martensite lathboundaries, and the cell walls or internal twins within the transformed regions). It is seenthat the solid-solution hardening effect of nickel accounts for about three-quarters of theoverall strength of Fe-Ni lath martensite containing 20%Ni. Figure 2.13 also showsthat changes in <r8 appear to correlate with the type of transformation substructure. Asthe transformation substructure changes at 4%Ni from a randomly arranged low-densityarray of dislocations to a much higher density of dislocations and a lath structure,<js increases sharply. However, there is no discontinuity in yield strength when the sub-

13

Oil U I" l 1>\1 U M N t : H l l . H - S I K l A t i i l l S 1 T 1 1 S

nk

Annealed Ferrite_

4S0

400:

3 5 0 ^

300S

2 S 0 5

76 73Co CONTENT at.%

SO

I it;. 2.14. — Transl'ormalinn sub-SII'UL'UIIV strengthening cll'ocisin 1 e-C'o alloys. After A C .S s u i v and I:..K. l i n v |.\V.>'|

structure changes from a lath- to a twinned-martensile type. As regards Fe-Co binaryalloys. Figure 2.14 shows that progressive strengthening occurs as the cobalt content isincreased, culminating in a 44",, increase in O.I "„ offset yield strength due to the formationof a lath mariensite substructure.

In Fe-C iillo\s with low carbon contents, up to about 0.05",,. most of the strength r r themariensite seems to deri\e from the high density of dislocations, which is probably about10i; to 10*; per cm :. In alloys with higher carbon contents, it seems likely that thestrengthening due to carbon exceeds by far that due to the substructure [2.44]. Recentwork on Fe-Ni-C alloys [2.45] has shown that, at carbon levels higher tl.an 0.3 wt.",,, thelath (cubici martensites are significantly stronger because they deform by slip whereastwinned I tetragonal I martensites with the same carbon content deform by twinning. It wassuggested that deformation by twinning is suppressed in the high-carbon cubic martensiusbecause in these structures all twin variants inevitably carry a large fraction of the carbonatoms into high-energy, non-octahedral sites.

The work-hardening behaviour of lath martensites. the temperature- and strain-ratedependences of their flow stress, and the derived values of the activation volumes andenergies for deformation, are all reasonably consistent with the behaviour observed iniron or low-alloy steels [2.46]. Furthermore, their Pelch slope values, based on the priorausteniie grain diameters, are on the low side of the results reported for io\v-a)loy steels.Thus, the properties of lath martensites can be considered as very similar to those ofcold-worked iron, as was proposed earlier [2.47]. One striking difference in propertiesbetween lath mariensiie and cold-worked iron is that the ductility of the former isreasonably good, as is exemplified by the Fp-l8Ni martensite even when tested in liquidhydrogen. Studies of the plastic deformation of an 18",,Ni binary [2.48] and severalFe-Ni-Cr ternary [2.49] lath martensites showed that the alloys all behave in a rathersimilar fashion. In the Fe-Ni-Cr ternary alloys, higher nickel levels considerably improvedthe impact properties at sub-zero temperatures Chromium was also slightly beneficial totoughness. Comparatively, pure iron cold-worked to a yield strength of 100,000 psi(700 MN m:) would have very little ductility. The reason for this difference may lie in thefact that the number of mobile dislocations is higher in lath martensites [2.48]. Theavailability of mobile dislocations in the structure would, in general, minimize cleavageand allow slip to take place. Thus, with regard to the maraging steels, one obviousadvantage of the Fe-Ni lath martensite matrix is that the flow stress is considerablyimproved as compared with that of pure iron, without any loss in ductility.

Heat treating lath martensite at temperatures up to 75O°F (400"C) produced fairlylarge increases in the elastic limit, and small increases in the yield and ultimate tensile

14


strengths [2.46]. Small decreases in residual mierostress [2.50] and electrical resistivity[2.5/] were also noted after heat treating Fe-25Ni maraging alloys. In the case of Fe-Ni-Cralloys, the increases in the elastic limit were inversely related to the A/« temperature of thealloys. The reasons for the observed changes in properties with tempering are not certain.Some may be due to the precipitation of trace amounts of carbon or nitrogen. However,it seems likely thai a recovery reaction to relieve the residual stresses generated by themartensitic transformation is primarily responsible [2.49].

There is also some evidence that, after tempering, lath martensite gives belter toughnessthan twinned martensite. The role of microtwins in lowering the toughness of Fe-Ni-Co-Csteels was demonstrated by comparing heavily twinned with untwinned martensile, andtwinned martensitc with bainite [2.J9]. At similar strength levels, the toughness of lowerbainite was found to be superior to that of heavily twinned martensite, but inferior tothat of untwinned martensite. Since, as stated • • Section 2.3.5, cobalt is not effective inreducing twinning in the higher-carbon steel* f :his type, the latter's toughness is notenhanced by cobalt additions; in fact, add.:ii<.;.i i t more than 4%Co was found to bedetrimental to the toughness of high-carbon .•.•.•.. 2.39].

In conclusion, this section on mechanical properties again emphasizes the importanceof lath martensite with respect to low-carbon maraging steels. In view of the general

£5& dependence of lath martensite formation on SFF. and Ms temperature, the desire to©P avoid both the formation of twinned martensite and the presence of untransformed

KJ austenite places rather definite limits upon the alloying additions that can be made to theFe-Ni base. From this point of view cobalt, which raises the Ms temperature of the basealloy, is very helpful in that it allows larger amounts of alloying elements to be addedwhile ensuring that the lath martensitic transformation remains possible. In addition,although cobalt does not affect the martensitic hardenability, it may possibly be used toincrease the martensitic hardness level by solid-solution strengthening.

2.3.7. Controlled Martensitic Transformation

The main feature of the controlled-transformation stainless steels is that the martensitetransformation temperature range is carefully controlled so that it is close to roomtemperature. The microstructure is also controlled so that it contains a small percentageof delta ferrite which aids carbide precipitation during the primary tempering operation.The advantage of these steels is that they have an austenitic structure which allows coldworking after cooling; the martensitic transformation is induced either by primarytempering which, through precipitation reactions, raises the Ms temperature of the residualaustenite, or else by refrigeration. The fact that the microstruciure is controlled meansthat the chemical composition is very critical, and consequently these steels are difficultto produce. Cobalt is a particularly useful element in such steels because it reduces thestability range of S ferrite without depressing the Ms temperature and it car. be used tobalance the effect of other alloy additions.

The general background to this type of steel was described by several authors [2.27, 2.52].It was shown that a 0.1C-17Cr-4Ni steel is a good base composition from which to developsuch steels. It is just on the borderline for S-ferrite formation and its martensitic trans-formation range lies just abOVe room temperature. I" order to modify the temperingcharacteristics, alloying elements such as molybdenum (one of the useful carbide-formingelements) and copper or aluminium (which are useful age-hardening elements) are added.When molybdenum is used, its ferrite-forming tendency is conveniently balanced bycobalt;'simultaneously, the chromium and nickel contents are reduced to allow additionof the increased amounts of cobalt and molybdenum. A typical composition of this typeis 0.06C-16Cr-3.5Ni-4Mo-6Co, which exhibits better impact properties than do the 12%Cr

i steels, as well as improved tempering resistance [2./].j"i 15

l"ORAL i -COM AININCi HK'iH-Sl RtNCiTH STEELS

500

as 1 1 6. 8 JO 20AGEING TIME, hours

<*) 60 80100 350

h

t//

I " -—

•

.

" • " — —

~ —

" - .

2% Co0% c i — _

AGEING TIME, hour*

at O.IC-rCr-JNiOMo steel, aged at 840 F (450 C). b) 0.05C-l7Cr-4Ni-4Cu steel aged at 84ITF (450°C).After K.J. IRVINE ei ai. [2.52]. After D. COUTSOURADIS 12.27}.

Initial condition : solution treated ai 1920'F Initial condition : solution treated at 1830°r (IGOCTC).(1050 O . and tempered at 1290°F (700°O. water quenched, and refrigerated at —112°F (—80'C).

Fig. 2.15. — Effect of cobalt on age-hardcnability of comrcr.jd-transformation steels.

In addition to helping to control the structure, cobalt has been found to enhance the agehardenability of the modified steels. This is shown in Figure 2.15 for a Cr-Ni-Mo and aCr-Ni-Cu steel. Since the general sequence of the tempering changes is identical to thatin Co-free steels, it has bean suggested that the effect of cobalt is almost entirely due tosolid-solution hardening [2.1].

2.4. Precipitation Strengthening

Precipitation reactions in steels are normally effected through tempering or ageing her.treatments. In carbon and alloy steels, it is usual to distinguish several stages in thetempering process (Fig. 2.16). The first, stress relief, corresponds to precipitation ofe carbide at temperatures of the order of 200-4003F (100-200°C); in actual fact, it is precededby carbon segregation or pre-precipitation clustering [2.36]. The second stage, whichtakes place between 400 and 600°F (200 and 300°C), involves the decomposition of retainedaustenke. Precipitation of cementite (FesC) in most carbon steels tempered between500 and 13OO'JF (250 and 700°C) is known as the third stage; at temperatures up to 750°F(400=C), Fe3C precipitates in a Widmanstatten structure, but above this it progressivelytransforms to spheroidal precipitation. Substitution of more stable alloy carbides forcementite at 900 to 1300°F (500 to 700°C) constitutes the fourth stage of tempering.Finally, the fifth stage involves the precipitation of intermetallic compounds. Beforediscussing the carbide and intermetallic-compound precipitation reactions in greaterdetail, the mechanisms which account for the strengthening role of precipitates will firstbe briefly reviewed.

2.4.1. Mechanisms

In discussing the various mechanisms proposed for precipitation or age hardening [2.54,2.5S], a distinction should be made between Guinier-Preston (GP) zone-type precipitatesand ordinary precipitates. The former may be regarded as a compositional change in thesolid solution, with or without a structural change. If there is a structural change, then it isa perturbation of the structure of the solid solution. An ordinary precipitate may becoherent, semi-coherent, or incoherent with the matrix. Actually all types of precipitate

16


Singe V

Stage IV

Swge III

Stoge II

Stage 1

Fig. 2.16.-

-59-TEMPERING TEMPERATURE'C

• i tjformation and growth •of Intennttoilic compounds^,

low-carbon marwmltt + £-r-ferrlte + cementite |

Formation ana* growthof alloy carbldttî

w in m mTEMPERING TEMPERATURE'F

- Sequence of reactions in carbon and alloy steels during tempering.After A. KASAK el al. [2.53].

may produce hardening, but GP zones and ordinary precipitates with some degree ofcoherency are more effective in this respect.

The problem which must be considered in order to understand precipitation hardeningis the interaction between a dislocation moving on a slip plane and a field of particles.The ways in which trie dislocation can pass beyond these obstacles are the following :

(i) The particles, if weak and closely spaced, may be sheared or fractured (Ansell andLenel's model). In this model detectable plastic flow would occur only when the particlesare being sheared or broken by the passing dislocations. Dislocations pile up againstsecond-phase particles and the latter rupture whenever the accumulated stress is largeenough. Thus, according to the final form of Ansell and Lenel's model [2.56], the yieldstrength of a dispersion-hardened material is given by :

7 P*<T = (To

(0.82(2.2)

where (JO is the matrix strength, C is a constant, and G", r and / are the shear modulus,the mean radius and the volume fraction of the particles, respectively.

(ii) The dislocation bends between the particles, leaving a dislocation ring about eachparticle (Orowan mechanism). In this model, the plastic strain results from the expansionof the dislocation loops surrounding the particles which intersect the glide plane. Theinitial shear yield strength is given by the modified Orowan relationship [2.54] :

. Gb 2In

2bw i t h q> = — (2.3)

where a0 and G are respectively the initial yield strength and shear modulus of the matrix,b is the Burgers vector of the dislocation, r and d are respectively the mean radius of theparticles and the mean planar interparticle spacing, and v is Poisson's ratio.

(iii) The dislocation line bypasses the particle by cross slip, leaving dislocation segmentsbehind (Hirsch model).

Several other theories, interpreting strengthening in terms of elastic misfit stresses,precipitate/matrix elastic moduli changes or increase in particle surface, have also beenproposed to account for the yield stress of precipitation-hardenable alloys [2.55, 2.56].

17

COBAU-l/ONl AIMNti IlKiHSTRFNGTH STKRS

800-

TEMPERATURE.t300 400 530

600 BOOTEMPERATURE.*F

Fig. 2.17.— Effect of tempering temperature (time:I hour) on hardness of Fe-C marlensites.After G.R. SPEICH and W.C. LESLIE [2.36].

TEMPEHINS TEMPERMURE, »C100 200 300 400 500 600

quenched600 800 1000 1200

TEMPERING TEMPERWURE,''F.

Fig. 2.18. — Effect of a 3.3% cobalt addition onthe tempering resistance of an AISI 4340-type steel.

After V.K.. CHANDHOK ei al. [2.59].

2.4.2. Carbide Precipitation

Many of the important properties of steels are affected by the precipitation of alloycarbides during heat treatment, which can produce marked secondary hardening.Figure 2.17 summarizes the complete process of tempering in Fe-C martensites and thecorresponding hardness variations. Alloying elements such as molybdenum and vanadiumare widely used because of the beneficial effects obtained by precipitation of their alloycarbides. Cobalt does not form carbides but may affect carbide precipitation indirectlyby preventing recovery of the dislocation substructure during tempering; this providesmore nucleation sites and a finer dispersion of the dislocation-nucleated carbides. Thesefeatures will be discussed further in Chapter 3.

The production of GP zones in a-iron also affords a means of strengthening steels [2.57,2.58], but here both substitutional and interstitial solute atoms are essential. In ironcontaining 2 to 5 wt. % Mo and about 0.2 wt. % N, GP zones have been shown to developat 840 to 1110°F (450 to 6OO'JC), the precipitation sequence being as follows [2.57] :(i) formation of coherent GP zones on |I00j ferrite matrix planes; (ii) formation of apartly coherent intermediate metastable phase, TJ' ; (iii) precipitation of a stable incoherentphase, TrFe.,Mo3N. The material thus produced is very hard, the yield strength of theiron reaching about one-half of its theoretical maximum. Moreover, the GP zones areunusually stable; they persist, and the material remains hard, even after heat treating forseveral hours at temperatures approaching 1290T (700°C). The production of GP zonesin ferrite is not limited to the Fe-Mo-N system; it has also been shown to occur in Fe-Mn-N,Fe-Cr-N and Fe-Mo-C alloys [2.57]. It seems that formation of substitutional-interstitialsolute-aiom zones must precede the homogeneous precipitation of alloy-element nitridesand cait-ides in most systems. Normal quenching plus ageing treatments promoteheterogeneous precipitation so that zone formation is seldom observed in practice.

1 STRENGTHENING MECHANISMS IN HIGH-STRENGTH STEELS

Strengthening by carbide precipitation is of importance in high-strength low-alloy steels,carbide-strengthened sieels (Chapters 3 and 4), and carbon-containing stainless steels(Chapters 8 and 9). A full discussion of the latter two types of steel will be found in therelevant chapters. As regards high-strength low-alloy steels, they are used in the martensiticcondition, after some light tempering to obtain adequate ductility. In actual fact, temperingis often performed at relatively low temperatures (— 390 F. i.e., 200 C). and results in(he formation of metasiable precipitates. In order to reduce the loss of hardness whichaccompanies this treatment, use can be made of a non-carbide-forming element such assilicon, which has quite a marked effect in stabilizing the precipitates, and hence inimproving the tempering resistance of the steels. Cobalt, even in small amounts, seemsto have the same effect on the tempering characteristics if the steel contains carbide-forming elements such as molybdenum or tungsten [2./]. The role of cobalt here wouldbe to impede the growth and coalescence of the precipitated particles by concentratingin the surrounding matrix, as is the case in high-speed steels. This effect has been utilizedin commercial steels. Figure 2.18 shows that addition of 3.3 %Co to an AISI 4340-typealloy steel (0.4C-0.6Mn-0.3Si-0.8Cr-0.25Mo-1.8Ni) has a significant influence on thetempering resistance. Similarly. Co-modified 4137 steel (0.39C-0.7Mn-lSi-l.lCr-0.25Mo-0.15V-lCo) includes both silicon and cobalt to modify the tempering behaviour of themartensite by shifting FeiC carbide formation to higher temperatures [2.60]. This steelhas adequate strength and extremely good resistance to notched fracture [2.60, 2.6/].

2.4.3. Precipitation of Intennetallic Compounds

Strengthening due to the "• fifth " stage of tempering, i.e., that corresponding to theformation of intermetallic compounds, is used in numerous steels including the maragingtypes. There are, at present, two families of high-strength maraging steels grouping,respectively, the non-stainless compositions, which typically contain 18%Ni (see Chapters 5.6 and 7), and the stainless compositions, which contain 10 to 15% chromium (seeChapters 8 and 9). Both families of steels contain fairly large cobalt and molybdenumadditions. Depending on the composition of the steel, hardening results from theprecipitation of both carbides and intermetallic compounds, or of intermetallic compoundsalone. After quenching, ageing at a relatively low temperature leads to a sharp hardnessincrease, the magnitude of which is related to the steel's total alloying element content.

As will be shown in Chapter 5 (Section 5.3.4), two precipitation processes are operativeon ageing Fe-Ni-(Co)-Mo maraging steels. The first takes place within the martensiticmatrix and is predominant for ageing temperatures below 840°F (450°C), while the secondoccurs preferentially on dislocations and is predominant at higher ageing temperatures.The first mainly involves the formation of coherent ordered precipitates, whereas thesecond leads to the precipitation of intermetallic compounds which have been identifiedas Ni3Mo, Fe2Mo, ti-FeMo and Ni3Ti. In these steels, precipitation of the stable Ni^Mo,FeiMo and tr-FeMo compounds is probably preceded by the pre-precipitation ofmetastable zones.

The yield strength of unaged maraging steels is typically of the order of 100,000 psi(700 MN/m2>. After ageing, it ranges from 200,000 to 300,000 psi (1400 to 2100 MN/m2).so that age-hardening improves this property by 100,000 to 200,000 psi. Severalinvestigators have pointed out that precipitation strengthening of the magnitude observedfor maraging steels can be accounted for quite reasonably by Orowan's model involvingdislocation motion between the precipitated particles [2.62, 2.63]. In fact, after ageing,the precipitate particles range in size from 100 to 500 A and are quite uniformly distributedin the matrix, with an average interparticle spacing of about 300-500 A; the correspondingstrength increment derived using Eqn. 2.3 is of the above-mentioned magnitude. Further-

U H M \IMMi HICH-SIKHNGni STKl-LS

more, it «.as observed [2.62] thai, after straining by 1 to 2",,, the structure of the l8Ni(25O)steel shows tangles of dislocations around the precipitated particles, and also instancesof dislocations bowing between the particles. Finally, the initially high work-hardeningrate of the steels is. in general, more consistent with the Orowan mechanism. On theother hand, some authors [2.6-i, 2.65] have suggested that Ansell and Lenel's model mightbe more plausible. In this case, the high strength in maraging steels would be due noionly to the very fine interparticle spacing, but also to the very high shear strength of theprecipitate particles. Observations of particle shearing lend support to this model [2.65].Finally, analysis of the temperature dependence of the flow stress [2.66] tends to showthat the effective shear stress varies as thai of mild steel and that the major strengtheningis due to long-range internal stresses such as would be developed by fine precipitateparticles.

One of the most interesting strengthening effects in maraging steels is that due to thecombination of cobalt and molybdenum [2.46]. It has been shown that the age hardeningobtained in Fe-Ni. Fe-Cr and Fe-Ni-Cr steels when cobalt and molybdenum are presenttogether is much greater than the sum of the strength increments produced when theseelements are added individually. The role played by cobalt is not clear, since this elementdoes not significantly participate in the precipitation reactions; numerous explanationshave been proposed, as will be seen in Chapter 5.

2.5. Strengthening by Thennomechanicai Treatment

The properties of alioys can be controlled by thermomechanical treatment (TMT), i.e.,by the introduction of plastic straining into the heat-treatment cycle. This leads tostrengthening by some of the mechanisms operative in all metals : solid-solutionstrengthening, grain-boundary and interface effects, dispersion strengthening, and strainhardening. The influence of TMT is particularly strong in steels, where the phase changefrom austenite to ferrite (or martensite) on cooling and the presence of carbon, which has agreater solubility in the aus'.enite and therefore precipitates in the ferrite, combine tomaximize the effects.

In recent review papers, TMT's were classified according to the position of the deformationin the heat-treitment cycle, i.e.. the structure that is deformed and the final structure

STABLE AUSTENITE

Critics! Temperature

VStart of Isothermal

Transformation

Room Temperature

Fig. 2.19.—Schematic time-temperature-transforma-tion diagram showing thermomechanica! treatments.

After E.B. KULA [2.67].

20


that is formed. These treatments are shown in Figure 2.19 on a conventional isothermaltransformation diagram for steel. The TMT classification is then as follows :

Class I : the austenitc is deformed before transformation and the martensite forms fromthis strain-hardened austenite. Ausforming, ausworking, ausrolling and hot-cold workingall belong lo this class.

Class II : the austenite is deformed as it is undergoing isothermal decomposition, i.e..the deformation is carried out at temperatures in the vicinity of A/,. The martensite formsduring the deformation of the steel in a metastable condition. Most frequently thistreatment is applied to stainless steels, which have MK temperatures slightly below roomtemperature.

Class III : the deformation is carried out after the auslenite transformation and may befollowed by reageing. These treatments correspond to the strain ageing of austenite trans-formation products (martensite, tempered martensite, bainite or pearlite). Various termshave been coined to designate such treatments : flow tempering, strain tempering,marstraining, and tempforming.

The above classification is useful because within one group similar strengthening mechanismsare operative. The Class II TMT's are probably the least complex. In the low-carbonmetastable austenitic stainless steels to which this treatment is generally applied, thestrength is determined by the relative amounts of martensite and austenite in the structure,as well as the magnitude of the work hardening of the martensite and retained austenite.For Class I treatments, part of the strengthening is caused by the structural refinementof the austenite and the resulting martensite, and also by the presence of defects in themartensite, which have been inherited from the strain-hardened austenite through thephase transformation. The major part of the strength improvement is associated withthe finer carbide dispersion which, while obviously leading to dispersion strengthening,results essentially in a higher dislocation density in the martensite. For Class III treatments,most of the strength increase arises during the work hardening, and. the rest duringsubsequent reageing. The carbides in the martensite increase the rate of work hardeningof the martensite and hence the dislocation density. During reageing, some dissolutionof the carbides may occur, leading to Cottrell locking and stress-induced ordering aroundthe dislocations, and ultimately to a reprecipitation process yielding a finer carbidedispersion. It was also found that plastic straining after initial ageing, followed by afurther ageing treatment, can produce marked changes in structure and properties; inparticular, prior plastic straining can considerably shorten the time cycle for heattreatment [2.68].

Many investigations have been made of the effects of TMT on carbide-strengthened9Ni-4Co steels (Chapter 4), Ni-Co-Mo maraging steels (Chapters 6 and 7) and stainlessmaraging steels (Chapters 8 and 9). The strengthening response of the 9Ni-4Co steels toTMT Was evaluated with particular emphasis on fracture toughness (c/. Section 4.2.1) :as regards the conventionally heat-treated steels, both strain tampering and ausformingof martensite extend the strength range of these steels and produce about the sametoughness level for a given yield strength; in the same way, strain tempering the high-carbon grades in the lower bainitic condition increases their tensile strength whileretaining acceptable toughness. As regards 18%Ni maraging steels, most of the workhas been directed towards determining the effects of applying TMT before maraging(c/. Section 6.2.1) : it was found that ausforming produced only very minor improvementsin the'final strength; on the other hand, marstraining gave more substantial strengthincreases. Ausforming and strain-ageing treatments have also been used on AFC-77, acobalt-containing high-strength stainless steel, in order to improve the compromise betweenstrength and toughness (cf. Sections 8.4 and 9.2.1).

21

( . ' O H M 1 - u A I \ | \ l * - ( , H K . I I S I R t N t i l H K U I H S

3. CARBIDE-STRENGTHENED STEELS — PHYSICAL METALLURGY

3 1. Introduction

Carbide-strengthened high-strength steels constiiule a major alternative to maraging>ieels W». Chapters 5, t> and '), whenever the use of the latter is not absolutely essential.The development of vrohalt-eomaining grades such as the HP 9-4-X and the I0Ni-8Co-Cr-Mo steels resulted from the need to combine high yield strength with good toughnessand adequate weldubility. The HP 9-4-X sieels contain *>u-oNi. 4°0Co, small amountsof the carbide-forming elements chromium, molybdenum and vanadium, and carbonin contents ranging from 0.20 to 0.45 "o. depending on the grade. They were developedby Republic Stee! Corporation [3.1 to 3 3]: iheir compositions are given in Table 3.1.

The lONi-sCo-Cr-Mo steel also presented in this table was developed by United StatesSteel Corporation with the purpose of providing a structural steel for the fabrication oflarge pressure \essels and hydrospace vehicles which incorporate weldments of heavysections. Steels of this type were initially considered for use as filler metals capable ofpro\iding yield strengths of 170,000-200.000 psi (1200-1400 MN/m-) combined withsatisfactory Charpy V-notch impact strength (more than 60 ft-lb at Of. i.e., 80 J at— 18 O [3.4]. Subsequently this type of steel was evaluated as a base metal [3.5. 3.6].The early developmental work on these Ni-Co-Cr-Mo steels, which pertained not onlyto 10°uNi grades, but also to some containing 5",,Ni [3.6. 3.7], was based on the conceptof combining the advantages of standard carbon martensite with those of maraging incompositions designed to exhibit a " dual strengthening " effect [3.8]. However, laterinvestigations [3.V] showed thai strengthening is entirely due to carbide precipitationduring tempering, though according to recent, as yet unpublished, work [3.10] the contri-bution of intermetallic-compound precipitation to strengthening would be significant formaterial in the overaged condition.

The microstruciural Features of these sieels will be examined in the following sections, aswell as the way in which they are affected by heat treatment; their relationship with thebasic mechanical properties will also be reviewed. Although some of the steels listed inTable 3.1 are no longer commercially available (this is the case for HP 9-4-45, supersededby certain lower-alloy steels, and HP 9-4-25, discarded in favour of HP 9-4-30 orHP 9-4-20), they will not be excluded from this and the following chapter, in view of thevaluable information which their development and study has generated.

TABLE 3.1. — COMPOSITION RANGES OF COMMERCIAL CARBIDE-STRENGTHENED STEELS <«t.%. bal Fe)

Steel designation

HP 9-4-45*

HP 9-4-30

HP 9-4-25 *

HP 9 4-20

lONi-Co-Cr-Mo

C

0.42-0.48

0.29-0.34

0.24-0.30

0.16-0.23

0.12

Mn

0.10-0.25

0.15-0.35

0.10-0.35

0.10-0.35

0.004

Si

0.10max.

0.10max.

OJOmax.

0.20max.

0.11

P

0.01max.

0.01max.

0.01max.

0.01max.

0.001

S

0.01max.

0.01max.

0.01max.

0.01max.

0.006

Ni

7.0-8.5

7.0-8.0

7.5-9.0

8.5-9.5

10.0

Co

3.5-4.5

4.25-4.75

3.5-4.5

4.25-4.75

7.8

Mo

0.2-0.35

0.9-1.1

0.35-0.55

0.9-1.1

2.0

Cr

0.2-0.35

0.9-1.1

0.35-0.55

0.65-0.85

1.0

V

0.06-0.12

0.06-0.12

0.06-0.12

0.06-0.12

—

Al

—

—

—

—

0.02

Year

1962

1962

1962

1966

1965

No longer commercially available.

22

i. CAKIilDE-STRF.NGTHF.NED STEELS - PHYSICAL METALLURGY

3.2. Continuous Cooling and Isothermal Transformations

Both continuous cooling and isothermal transformation curves are available for steelsof the HP 9-4-X type; this is not the case, however, for the IONi-8Co-Cr-Mo steel.

3.2.1. Continuous Cooling Transformation (CCT) Curves

The CCT curves for four 9",,Ni steels are presented in Figure 3.1. Steels Z-9 and Z-9-4are experimental steels which contain, in addition to nickel, 0.3 ",,C and. respectively.0 and 4",,Co; the remaining two steels are industrial grades.

The CCT diagram for Z-9 shows, in addition to the martensitic transformation whichstarts at about 445 F (230 C), a broad bainitic transformation range beginning at about930 F (500 C), with a critical cooling rate of about 270 F/min (150 C min). No trans-formation to pro-bainitic ferrite was observed, even at the slowest cooling rate used (about4.5'F/min, i.e.. 2.5 C/min). On decreasing the cooling rate, a considerable, thoughprogressive decrease in hardness is evident, and the residual austenite content is found toincrease from 0% in the martensitic range to 15°o in the bainitic field.

Comparison of this diagram with that for the experimental Z-9-4 steel shows that cobaltraises the transformation points, both on heating (Af and Af) and on cooling (Ms andfls); furthermore, the critical cooling rate has risen sharply to 1800 F min (1000=C min).clearly illustrating that cobalt exerts an a-field broadening effect and increases the nucleationrates. It is also apparent that the residual austenite content in the bainitic range is lowerfor the 9Ni-4Co than for the 9Ni steel.

The reason for the difference between the 9Ni-4Co laboratory heats and the industrialsteels studied lies essentially in the latter having a lower nickel content (7.5 "„) and contain-ing the carbide-forming elements chromium, molybdenum and vanadium. The diagram

i COOLING RATE BETWEEN 1470 AND 930'F.T/Tninuie

600

MO

ir 200'° 51I

IB?

• + P H U F l-M-'t ! ' ' - - 4TTTffli

» 10 UTIME FOR COOLING BETWEEN 1A70 AND 93O'F(B0O*C AND 50O'C),minutK

Fig. 3.1. — Continuous cooling transformation diagrams for two experiment?.!and two industrial 9%Ni steels. After D. COUTSOURADIS ct ol. [3.JI].

SteelAustenizing temperature

Time at austenizing temneratureA,A,

0.3C-9Ni (2-9) 0.3C-9Ni-4Co (Z-9-4) HP 9-4-2015IO°F(82O°C) 1455°F (790°C) 1525°F (S30°C)

30 min1204°F (651°C)1114oF(602°C)

t5 min 30 min

HP 9-4-30I525°F(830°C)

30 min1330oF(722°C) 1567°F (853°C) 1546°F (841°C)1124°F (607°C) 1198°F (648°C) 1231°F (666°C)

23

U i | l \ l l - i O M M M S i . l l l i i l i M R l M i l l I M M I S

1 1 0 ' K> 2

TIME.minutes

Fin. 3.2. — Isothermal transformation diagrams for HP9-4-20 after R.T. AULT [3.12]. HP9-4-25after G.D. Rits ami S.W. POOLF [.*..?]. "and HP 9-4-45 after T.P. GKOENFVELD tri al. [3.131

for the HP 9-4-30 steel (0.3°,,C) shows that the carbide-forming elements lower both thecritical cooling rats Co 55 Fmin. i.e., 30:C/min) and 6.,-. On the other hand. A/* is onlyslightly raised, the effect of adding the carbide-forming elements being more than counter-balanced by that of decreasing the nickel content. Finally, the cooling rate does not affecteither the hardness or the residual austenite content: the latter, which is of the order of5",,. ir not decreased to any appreciable extent by holding for 2 hours at —IIODF (—80 C)after quenching. The lower carbon content of HP 9-4-20 leads, as expected, to an increasein the \tx and Bs temperatures, as well as in the critical cooling rate, which is raised to600 F,min (33OC. min). The residua! austenite content increases on passing from themartensiiic to the bainitic range: in the latter, the hardness decreases slightly.

3.2.2. Isothermal Transformation Curves

Tentative TTT diagrams for three HP 9-4-X grades are presented in Figure 3.2. Thesediagrams confirm that the steels are characterized by a well-defined bainitic range,particularly for HP 9-4-45.

3.3. Bainitic Transformation Structures

The high-strength steels considered here are generally used in the quenched-and-temperedcondition. However, as will be seen further on, heat .treating the higher-carbon grades,particularly HP 9-4-45, to form lower bainite isothermally results in an optimum combi-nation of strength and toughness, which is not the case for the lower-carbon grades orthe 10Ni-8Co-Cr-Mo steel. Furthermore, bainitic structures may develop in heavy-sectionweldments which are not subjected to post-welding heat treatment. It is thus appropriateto discuss these structures here; this will be done essentially in relation to the HP 9-4-Xsteels. The bainite structures formed both on continuous cooling and isothermal holdingwill be described, and their effect on strength and toughness examined.

3.3.1. Bainites Formed on Continuous Cooling

Austenized 0.3C-9N1 steels subjected to continuous cooling within the bainitic trans-formation range exhibit a microstructure that is typical of bainitic acicular ferrite [3.11].As shown in Figure 3.3a, the structure is rather coarse and is composed essentially oflarge areas of bainitic ferrite and white, less deeply etched zones composed of an intimatemixture of martensite and residual austenite. Although the martensite is clearly visible,

24


:^ rci

. -v «y

4\ + ** \

* > • >,.

<j(0 K 9 N i l 7 91 - 2 2 F m i n d 2 L mm) A) 0 1C 9 N i - 4 C o ( Z « 4 ) — 4 s F m m ( 2 i C miniAu ii-nizjiion 1^10 K R 2 0 O - I O m m WOO Au u n u a u o n 14^5 F<790 C> lMtun 3000

Fig. 3.3. — Replica electron micrographs of experimental 9%Ni steels aftercooling within the bainitic range. After D. COUTSOURADIS er al. [3.11].

there are practically no carbides. Addition of 4%Co results in a much finer structure(Fig. 3.36). Since the prior austenite grain size was the same in both steels, it is clear thatcobalt favours nucleation of the bainite and slows down its growth during transformation.

The bainitic microstructures of the industrial HP 9-4-20 and HP 9-4-30 steels differconsiderably from those of the steels without carbide-forming elements. Following slowcooling in the bainitic range, HP 9-4-20 presents a bainite-ferrite-type microstructure(Fig. 3.4a) as did the experimental 9Ni-4Co steel, but this is now accompanied by a finecarbide precipitation within the ferrttic regions, and the size of the austenite-martensiteislands is reduced. Higher cooling rates (of the order of 45°F/min, i.e., 25°C/min) lead, asis shown in Figure 3.4£>, to the formation of acicular ferrite which recalls martensite and inwhich cementite discs lying at an angle of 60° to the axis of the ferrite needles are observed;moreover, the crystallographic orientation relationships between the ferrite and thecementite are identical to those prevailing in tempered martensite. This morphology istypical of lower isothermal bainites. It should be noted that the carbides in the athermal

a) Replica— 4.5oF/min (2.5°C/min). x 3000 b) Thin foils — 45°F/min (25°C/min). X40.000Fig. 3.4. — Electron micrographs of HP9-4-20 after cooling within the bainitic range.Austenization : 1525°F (83O°C) - 30 min. After D. COUTSOURADIS et al. 13.11).

25

M M V . ( I l l , H S I R f N C . r u S TLLLS

„•) Replica - 4.5CF min (2,5=C;minK -3000 b) Thin foil — 4.5°F/min (2.53C/minl. v 10.000Fia. 3.5. — Electron micrographs of HP 9-4-30 after cooling within the bainitic range.Austenization : 15^5 F (S30C) • 30 min. After D. COUTSOURADIS et al. [3.11\.

bainites formeU at lower cooling rates also occur as discs aligned in a single direction. Inthis respect, these structures also resemble that of isothermal lower bainite.

After cooling at high rates in the bainitic range, the structure of HP 9-4-30 also containsappreciable amounts of martensite. Following slower cooling (Fig. 3.5a), precipitationof very fine carbides is observe \ within an acicular ferriie. These findings were confirmedfrom examination of thin foils; in particular, carbide particles oriented in one directionare once more evident after slow cooling (Fig. 3.5ft).

In summary, it can be said that continuuus cooling of HP 9-4-20 and HP 9-4-30 leads tostructures similar to that of lower bainite. The morphology of ihe carbides changesaccording to whether they precipitate within martensite or bainitic ferrite; this difference is

MEAH COOLING RATE BETWEEN U70 AND S30-F."F/minuH

nf raL io' iMEAN COOUNG RATE BETWEEN80OAND50O*C,*C/minute

io' sL SL

TEST TEMPERATURE,"':-120 -60 -IS 0_; 40 80 120

-200TIME FOR COOLING BETWEEN 1470ANDgM'FleOOANDSOQ'CKminules

- 1 0 0 - - 0 : • . - . . -•••-• 100--, V 2 0 0

TEST TEMPERATUBE,°F

Fig. 3.6. — Effect of cooling rate on the mechanical properties ofHP 9-4-20 and HP 9-4-30. After D. COUTSOURADIS et al. [3.11].First row of data points corresponds to martensitic range, second andthird to bainitic range. Austenization : 1525=F (83O°Q - 30 min.

Fig. 3.7. — Temperature dependence of Charpy V-notch impactstrength of HP 9-4-20 and HP 9-4-30 after cooling within thebainitic range. After D. COUTSOURADIS et al. [3.1]].

Austenization : 1525°F (83O°C) - 30 min.

26

.1. CARBIDE-STRENGTHENED STEELS — PHYSICAL METALLURGY

probably due to the fact that in the latter case, the carbides precipitate behind the ferrite-austenite interface as the ferrite needles grow.

3.3.2. Mechanical Properties Associated with Continuous-Cooling Bainites

The effect of cooling rate within the bainitic field on the mechanical properties of HP 9-4-20and HP 9-4-30 is shown in Figure 3.6. To a first approximation, and with the exceptionof the ultimate tensile strength of HP 9-4-30, the properties vary moderately with coolingrate. It can be seen that the continuous-cooling bainites in these steels, which, as pointedout above, are comparable to isothermal lower bainites, behave simiiariy To st-lf-ternperedmartensite (cf. Section 3.5.1). As shown in Figure 3.7, which is a plot of impact strengthversus test temperature, the " massive " bainite (for definition of this term, see Chapter 2,Section 2.2.1) obtained in the HP 9-4-20 steel cooled at a rate of 4.5=F/min (2.5cC/min)is the only one which exhibits an appreciable decrease in impact strength at lowtemperatures.

3.3.3. Isothermal Bainite in HP 9-4-45

Isothermal bainite can be formed in HP 9-4-45, which has an Ms temperature of about465°F (240-C), by austenizing at 1500°F (8I5=C), hot quenching into a salt bath at 450to 75O°F (230 to 400°C), holding for up to 16 hours, and finally oil quenching [3.14]. Thebainitic reaction starts after 7.5 to 15 minutes at all transformation temperatures, andstops within two hours of isothermal holding, although this does not necessarily mean thatthe final structure consists of 100% bainite. For example, the end of the bainitic trans-formation at 750 and 700° F (400 and 370°C) corresponds to approximately 80 and 90%bainite, the remaining phases (upon quenching) being retained austenite (3 and 7 vol. %.respectively) and untempered martensite. Transformation temperatures of 65O-500cF(340-260°C) appear to give fully bainitic microstructures. Decreasing the reactiontemperature to 475CF (245°C) produces 98% bainite and 2% retained austenite. Thestructure generated by holding at 450°F (230°C), i.e. below the Ms temperature, is a morecomplex mixture of tempered martensite, bainite and retained austenite.

The microstructural features of the bainites thus formed vary with transformationtemperature. The bainite obtained between 450 and 600°F (230 and 315°C) exhibits thetypical lower-bainite morphology. It should be noted that this temperature range is thatwithin which cementite precipitates during tempering of martensite [3.15] and that, asalready mentioned, there is a great structural similarity between lower bainite andtempered martensite. The bainitic plate size increases as the reaction temperature israised. A marked transition occurs between 600 and 650°F (315 and 340°C). At 650°Fand above, the bainite consists of ferrite plates with strings of cementite that tend to lieparallel to the major direction of the ferrite, both in and between the bainitic plates.However, as the transformation temperature is raised further, the upper bainite becomescoarser and the carbides form mainly between the bainitic plates [3.14].

3.3.4. Mechanical Properties of Isothermal Bainite in HP 9-4-45

The strength and fracture toughness properties of the bainite in HP 9-4-45 are presentedin Figure 3.8 as a function of reaction temperature. The yield and ultimate tensile strengthsdecrease progressively as the reaction temperature is raised from 450 to 700°F (230 to370°C), with an accentuation in this trend for the yield strength over the 600 to 650°F(315 to 345°C) interval, within which the structure changes from lower to upper bainite.There is no retained austenite in the subcooled and tempered bainites formed at 500 to650°F (260 to 345°C), and less than 1 % in those formed at 450, 475 and 700°F (230, 245and 37O°C). A reduction in area of 35 to 45 % and an elongation of 7 to 8 % are observedfor all these reaction temperatures.

27

I OB-XL 1-CON I AININCi HlCHl-SHU NCi ITl SI I HI S

REACTION TEMPERMURE.'C

ZO ??5 30Q 325 350

H 7.-1 -- energy per unit area determined on pre-eraekud substandard-thickness Charny V-nolchspecimens I2.ih5 • • 0.394 • 0.08 in., i.e. 55 < \O • 2 mm>.A» -• critical stress intensity factor associate-! with the initiation of unstable fracture in slowbend lest on prccracked substandard-thickness Charpy V-nolch specimens (2,163 •• 0..VM- 0.2 in., i.e. 55 x 10 < 5 mm).

Fig. 3.S. — Effect of isothermal holdingtemperature on the mechanical propertiesof HP 9-4-45. After D. KAUSH et at. [3.14].

Austemzation : 1500F (815C) - 30min:bainite reaction time 6 h, followed by oilquenching, refrigeration in liquid nitrogen, anddouble tempering (1 - ! h) at 400 F (2O5°C).

x 10.000 x 40.000Fig. 3.9. — Thin-foil electron micrographs of HP 9-4-20 after martensitic quenching.After D. COUTSOURADIS el ai. [3.1i\. Austenization : I525°F (83OX) - 30min, O.Q.

The precracked impact energy (WjA) of the lower baintte increases as the reactiontemperature is raised from 450 to 600CF (230 to 315°C) (Fig. 3.8), but the transition fromlower to upper bainite is accompanied by a decrease in WjA as well as a reduction instrength; as the upper-bainite formation temperature is raised, WjA decreases further.The toughness behaviour of the upper bainite is also reflected by the lower values of thecritical stress intensity factor {Kv). It is thus evident that the lower bainites offer attractivecombinations of strength and toughness, whereas the properties of the upper bainitesare clearly disadvantageous. A reaction temperature of 500°F (260°C) gives a yield strengthof 205,000 psi (1415 MN/m*), a tensile strength of 250,000 psi (1725 MN/m2), and a Kq

value of 101,000 psiy'in. (110 MNm-J '2) [3.14].

3.4. Martensitic Transformation on Quenching

As stated in Chapter 2 (Section 2.3.5), the occurrence of either lath or twinned marten-site appears to be controlled, at least partly, by the carbon content of the steel. As amatter of fact, the high-nickel steels considered here, with the possible exception of thehigh-carbon grades (.-.g. HP9-4-45), exhibit a microstructure of dislocation-rich lathmartensite in the as-quenched condition [3.9, 3.11]. The higher Ms temperature due tothe cobalt addition results in a martensite which is virtually free from austenite, but inwhich some self-tempering has occurred. Thin-foil electron micrographs of HP 9-4-20after martensitic quenching are presented in Figure 3.9; it is seen that the cementiteparticles are arranged in a Widmanstatten stiucture typical of tempered martensite.

Increasing the austenizing temperature results in a larger austenite grain, and hence in agreater martensite lath size. As a result, the yield strength decreases, as is illustrated inFigure 3.10 for the 10Ni-8Co-Cr-Mo steel (the reverse trend observed above 1600°F (870°C)is due to the concurrent effects of more carbon going into solution and strengtheningthrough self-tempering) [3.16]. Double austenizing treatments have been shown to givebetter combinations of strength and toughness because they make it possible to optimizeboth the austenite grain size and the amount of dissolved carbon [3.7]. In the case of

28


5Ni-Cr-Mo steels with carbon contents between 0.15 and 0.25% and cobalt contentsfrom 0 to 8%, a double austenizing treatment of 1 hour at 1650°F (900°C) 4- 1 hour at1600T (870uC) raises the room-temperature Charpy V-notch impact strength by 3 to13 ft-lb (4 to 18 J) over that of the steel austenized simply for I hour at I500°F (8I5°C),both types of treatment being followed by tempering for 1 hour at IOOO°F (540°C). More-over, the beneficial effect of the double treatment was found to be independent of thecarbon and cobalt contents [3.7], Double austenizing treatments have also been recom-mended for welded assemblies of the lONi-Co-Cr-Mo steels [3.8]. Prior to welding, the1 in. (25 mm) plate used was fully heat-treated as follows : 1660°F (905'C) - 1.5 h, W.Q. +1500"F (815 C) - 1.5 h, W.Q., + 950°F (510"C) - 5 h, A.C.

3.5. Tempering Reactions

The subjects dealt with under this heading are the tempering reactions that occur in thevarious types of steels considered in the present chapter, and their effects on mechanicalproperties. For a general discussion of these reactions, the reader is referred to Chapter 2,Section 2.4.1.

3.5.1. HP 9-4-X Steels

In the as-quenched condition (cooling at the rate of 5400°F/min, i.e., 3000DC/min),HP 9-4-20 is essentially comprised of self-tempered acicular martensite. After temperingfor 2 hours at 570°F (300°C), areas of acicular ferrite containing abundant cementiteprecipitates can be distinguished (Fig. 3.11a). Tempering at higher temperatures doesnot significantly modify this structure, but the triaxiality of the cementite needles becomesmore marked; for tempering temperatures between 750 and 930°F (400 and 500°C), M7C3chromium carbides have been observed in addition to the cementite [3.1]]. In HP 9-4-25,only cementite was identified after tempering between 200 and 1100°F (100 and 600°O[3.15]. As-quenched HP 9-4-30 also presents a self-tempered martensitic structure madeof untwinned needles; however, the extent of self-tempering is less, on account of thelower Ms point of this steel as compared with HP 9-4-20. On tempering at 570, 750 and930°F (300, 400 and 500°C), abundant precipitation of very fine carbides, which do notappear to coalesce, occurs within the martensite (Fig. 3.lib). The following crystallo-

& 165

Sx5 160

Q 1S5

! 150

MM

AUSTENIZING TEMFERATURE,°C

800 650 900

\

1

• •

\

\

1 ez2

IIOO!

150O 1600

AUSTENIZING TEMPERATURE,°F

Fig. 3.10. — Effect of austenizing temper-ature on as-quenched yield strength of10Ni-8Co-Cr-Mo steel. After T.B. Coxand A.H. ROSENSTEIN [3.16\. Austenizingtime 1 h, followed by water quenching.

a) HP 9-4-20 — Thin foil.x 40,000

b) HP 9-4-30 — Replica.10.000

Fig. 3.11. — Electron micrographs of HP 9-4-20 and HP 9-4-30 afterquenching and tempering. After D. COUTSOURADIS et at. [3.11].Condition : 1525T (83O°C) - 30 rain, O.Q. + 570°F (300°C) - 2 h, A.C.

29

IOHM.1COM \1MNI, HH.H-STRFNGIH SI EELS

TEMPERING" TEMPERATURE, "C• 300 400 500

TEMPERING TEMPERATURE,°C300 <00 500

H2Q00

Z-9(0.3C-9Ni)Z-9-4(Ci,3C-9Ni-4Co)

quenched quenfhed TEMPERING TEMPHRATURE.'F

ut Experimental steels b) Industrial steelsA,,-, .„,„.,;„„ f z ' 9 : 15'0°F (82O'Ci - 30 min. O.O. Austeniza'tion : 152'."F <83O°C> - 30 min. O.O.

Fig. 3.12. — Effect of tempering temperature on the mechanical properties oftwo experimental and two industrial 9" oNi steels quenched within the marten-sitic range and tempered for 2 hours. After. D. COLTSOURADIS et al. [3.11).

graphic orientation relationships between cementite and ferrite in these steels have beenestablished from electron microdifiraction examination of thin films :

(001)Fe)C//(2ll), [100]FejC//[0Tl]3 [010]FejC//[]TTL. / "They are in excellent agreement with those reported in the literature. In HP 9-4-45, theprecipitate formed on tempering at temperatures up to 35O°F (!75°C) ;s the hexagonalz carbide; the carbide formed on tempering between 400 and 1100"F (200 and 60C°C) iscementite [3. IS].

The tempering response of 0.3C-9Ni and 0.3C-9Ni-4Co«steels is compared in Figure 3.12a.The properties given, especially Ihe reductions in area, elongations and impact strengths,are very similar for both steels and exhibit the same dependence on tempering temperature.However, the ultimate tensile and yield strengths of the 0.3C-9Ni-4Co steel do not fallwhen the tempering temperature is raised from 750 to 930°F (400 to 500oC); this alreadyprovides some indication of the relative insensitivity of the steel to heat treatment. Thisis even more apparent in the case of the low- to medium-carbon industrial grades(Fig. 3.126), whose properties hardly change with tempering temperature between 570and 930JF (300 and 5OO'JC). It is seen, in particular,-1 that the yield strengths of the 0.3and 0.2%C steels remain around 200,000 and 170,000 psi (1400 and 1200 MN/m^),respectively.

The industrial heat treatment of HP 9-4-20 consists in quenching followed by singletempering for 4 hours at 1O25°F (550aC); a similar treatment is recommended for HP 9-4-30.

30

i. CARBIDE-STRENGTHENED STEELS - PHYSICAL METALLURGY

except that the steel is cooled clown to —110'F (—80°C) after quenching and then doubletempered for 2 • 2 h at about IO00F (540'C). The refrigeration step for HP 9-4-30produces an increase in yield strength of about 15,000 psi (100 MN/m-) [3.17]. Thestability of the industrial grades with respect to tempering temperature is no doubt related,on the one hand, to the presence of cobalt which slows down the tempering process and.on the other, to the secondary hardening that results from the precipitation of secondarycarbides.

3.5.2. 5Ni-Cr-Mo Steels

As staled in the introduction, part of the developmental work on lONi-Co-Cr-Mo steelswas devoted to several 5%Ni grades. This included the determination of the effect oftempering temperature on the yield strength and Charpy V-notch impact strength of0.25C-5Ni-1.5Cr-0.5Mo steels containing increasing amounts of cobalt [3.6, 3.7]. In theas-quenched condition, the microstructure of these steels is once more typical of lowerbainite or self-tempered martensite. Tempering in the 600 to 900°F (315 to 480cC) rangeresults in the formation of FejC grain-boundary films, which are responsible for the lowtoughnesses measured in this range (Fig. 3.13). After tempering at I000°F (540~C), thegrain-boundary Fe3C carbides are discontinuous as a result of spheroidization, hence theincrease in toughness; at this tempering temperature the steels .are reported to containboth equilibrium-type M6C and FejC carbides [3.7]. The microstructure of the cobalt-modified grades after quenching and tempering at 10005F (540°C) is characterized by thepresence of a very fine carbide precipitate in the matrix. This is believed to be responsible

100TEMPERING TEMPERATURE. °C

20D 300 400 500 700

200 400 500 800 1000TEMPERING TEMPERATURE,0 F

1200 1400

Fig. 3.13. — Effect of cobalt content on tempering responseof 5Ni-Cr-Mo steels. After L.F. PORTER el al. [3.6].Base composition : O.25C, 0.75Mn, 5Ni, 1.5Cr, 0.5Mo.

Austenization : 1500°F (815°C) - 1 h, W.Q. Tempering time : 5 hours.

31

(.•OH-\LT-l'O\r.\lNINl. HKiH-STRKNlim STRF.LS

TEMPERING TEHPERATURS.'C

Y « VSTABLE — TRANSF.-J500

ON I ONCOOLING COOLING

6 m BOO 1D0OTEMPERING TEMPERATURE, °F

1000 1200 1(00

Fig. 3.14. — Effect of carbon content on tempering responseof lONi-Co-Cr-Mo steels. After G.R. SPEICH el al. [3.9].

Base composition : !ONi-8Co-2Cr-lMo. Austenization : 155O°F (845'Ci •1 h, W.Q. Tempering for 1 h, followed by water quenching.

for the higher yield strengths and possibly the decreased notch toughnesses of the cobalt-containing steels.

3.5.3. lONi-Co-Cr-Mo Steels

According to a recent study [3.10], both carbides and the intermetallic compound Ni3Moare liable to form in the !0Ni-8Co-Cr-Mo steel during heat treatment. However, thestrengthening contribution of Ni3Mo was only found to become significant on overageing,while the observations made during the initial stages merely confirmed the homogeneousprecipitation of carbides generally held to be responsible for the steel's strength. Thepresent section will therefore be devoted exclusively to the latter type of precipitate.

The reactions occurring in 10Ni-8Co-Cr-Mo steels with three different carbon contentsin relation to tempering temperature are summarized in Figure 3.14. The hardnessdecrease observed when the tempering temperature is raised from 400 to 800T (205 to425=C) is associated with the precipitation of M3C (where M = Fe, Cr, Mo) in aWidmanstatten structure. Between 800 and 1000°F (425 and 5^0°C), a secondary hardeningpeak is observed; this is related to the dissolution of M3C and the formation of a finedispersion of (Mo,Cr)2C needle-like carbides which nucleate on dislocations and growat the expense of M3C until the latter eventually disappears. From 1000 to 1200°F (540 to650°C), where the hardness drops sharply, the M2C carbides coarsen rapidly, the sub-structure is rapidly annealed out, and much of the austenite formed at the tempering

32

V CARUIDF.-STRF.NriTHF.NED STEELS — PHYSICAL METALLURGY

temperature is stable upon quenching to room temperature. At 1200 to 1300cF (650 to705 C), the hardness starts to increase again, because the auslenite formed upon temperingnow has a lower alloy content and transforms completely lo mariensite upon quenching.Finally, above I400F (760 C), complete austenization occurs and the hardness returnsto the as-quenched value [J.f].

The microstructures of the 0.09 and 0.I9C grades tempered for I hour at 950F (510 C)are shown in Figure 3.15. The plate-like M3C particles, the MjC carbides nucleated ondislocations, and the high dislocation density of the mariensitic matrix are clearly apparent.The progressive dissolution of M3C and coarsening of the M;C carbides as the temperingtemperature is raised are shown in Figure 3.16 for the 0.13C grade. The fine dispersionof dislocation-nucleated M^C carbides is enhanced by the presence of cobalt, whichincreases the dislocation density of the matrix. Figure 3.17 illustrates this effect in thecase of a 0.12C-10Ni-2Cr-l Mo steel. Retention of a high dislocation density in the cobalt-containing grades is in agreement with observations according to which cobalt raises therecrystallization temperature of 9%Ni steels [3.18].

150,000 a) 95O'JF (510"cl - 12h

150.000 b) 100O°F(54O°C)-12h 85,000 b\ 8"0Co y. 85.000

'/jFig. 3.15. — Effect of carbon content oncarbide precipitation in 10Ni-8Co-2Cr-lMosteels tempered for 1 hour at 950°F

JSKTC). After G.R. SPEICH et al. [3.9].irAustenization : 155O°F (845°C) - 1 h, W.Q.

Fig. 3.16. — Effect of temperingtemperature on carbide distributionin 0.13C-10Ni-8Co-2Cr-lMo steel.After G.R. SPEICH et al. [3.9].

Austenization : 155O°F (845°C) - 1 h, W.Q.

Fig. 3.17. — Effect of cobalt on recovery ofdislocation substructure in tempered 0.12C-10Ni-2Cr-lMo. After G.R. SPEICH et al. [3.9).

Condition : I55OCF (845°C) - 1 h.W.Q. +950°F(5103C)-12h.A.C.

33

1 . H U M V i O N I \ l M \ l i H R . H - S I K I N l H H S T I f l S

1000 1200 1400TEMPERING TEMPERATURE.'F

I K . ? . |S . HTeet of tempering temper-ature on austenite formation m H)"f,Ni-tceK. -\Her Ci.K. SPHIH el ill. [3-'-*\.•Vii-tcni/ation : 1550 K (S45 O - I h. W.Q.

Tempering time : 1 hour.

hiu 3.11). I tl'ect of individual alloyingelements on tempering response of 10",,Nisieelv After OR. SPHI ii ft at. \3.t\.AitMeni/ai-.nn : 1550 V (H45 Cl - ! h. W.I).

Tempering time ; 1 hour.

TEMPERING TEMPERATURE.'C100 200 300 «X) 500 600 700

ai2C-ONi-6C<J-2C|--iMo_

!0.12C-l0Ni-8Co-2Cr

ffl2C-10Ni-aCo

Q12C-10N1

DNi :

1300

1200

noo :

VWOg

3SCO Z

800

600

500

RT " 200 «X) 600 600 1000TEMPERING TEMPERATURE.'F

1200 VO0

3.5 4. Retained .-Ui.slcnite

Formation of austcnite during tempering and its tetention upon quenching from thetempering temperature have been studied in the case of a 0.12C-iONi-8Co-2Cr-!Mo and aO.I2C-IONi steel [_?.<>]. Figure 3.18 shows the amount of austenite retained at roomtemperature after tempering for 1 hour. No retained austenite is observed in either steelafter full austenization and quenching: ii only starts to occur at temperatures exceedingthe secondary hardening peak, in amounts which increase as the tempering temperature isincreased to 1200 F (650 C) and then decrease to zero at 1300=F (705"C). Prolongedexposure (up to 20 h) at 950 F (510 C) does not result in any austenite Eormation, whereasexposure at 1000 F (540 C) gives rise to a progressive increase in retained austenite content,with a parallel decrease in yield strength [3.9].

3.6. Effect of Alloying Clements on Tempering Response, Strength and Toughness

As a general introduction to this section. Figure 3.19 shows the tempering behaviour of alow-carbon Fe-IONi base and its modifications following successive additions of carbon,cobalt, chromium and molybdenum [3.9]. After quenching, the base alloy has a low yieldstrength which is retained on tempering at temperatures up to about I000T (540 C).Addition of 0.12",,C raises the as-quenched yield strength, but appreciable softening isobserved on tempering at temperatures above 400"F (205C). Addition of further elements(Cr. Co, Mo) does not significantly affect the as-quenched yield strength, but has a stronginfluence on this property after tempering. Cobalt and, to an even larger extent, chromiumincrease the steel's resistance to softening, whereas molybdenum contributes to secondaryhardening. These effects wiil be examined in greater detail in the following sub-sections.

34

f

.! ( AKBlOK-SIKKNCilHINI.U STEELS - PHYSICAL. METALLURGY

027. OFFSET YIELD STRENGTH, MN/m2

110Q 1200 1300

"150 160 170 180 190 2000.27. OFFSET YIELD STRENGTH,

Fig. 3.20. — EITect of carbon contenton notch toughness - yield strengthrelationship for IONi-8Co-2Cr-lMosteel. After G.R. SPEICH el at. [3.9].

- \ vniiation : 1550'F (S45C) - I h, W.QInnpering for indicated times at 950"F (510 C).

160,0.25 0.30 0.35 0.40

CARBON, wt.%0.30 0.35 0.40

CARSON, wt .V.

Fig. 3.21. — Effect of carbon content on mechanicalproperties of HP 9-4-X steels in tempered martensiticcondition. After J.S. PASCOVER and S.J. MATAS [3.2].Condition : austenization, oil quenching, refrigeration,and 2 - 2 h tempering at indicated temperatures.

3.6.1. Effect of Carbon

When the carbide-strengthened Ni-Co steels were developed, the effect of different carboncontents was evaluated from the start, since this element is obviously responsible for basichardening through carbide precipitation and interstitial solid-solution strengthening.This essential role of carbon was illustrated in Figures 3.14 and 3.19. Figure 3.20 showsthe toughness vs. yield strength relationship for 10Ni-8Co-Cr-Mo steels with differentcarbon contents (0.09, 0.13 and 0.19%) tempered for increasing times at 950°F (510°C).The graph exhibits C-curve behaviour because of the secondary hardening that occursat this temperature and of the simultaneous increase of both yield strength and toughnessthat results under some tempering conditions.

The effect of carbon on the strength and toughness of the HP 9-4-X steels is illustrated inFigure 3.21. The strength increases with increasing carbon content, whereas the toughnessdecreases; this effect is less marked for the higher-carbon grades. Increasing the carboncontent aiso decreases the ratio of notched tensile strength to yield strength, irrespectiveof the strength level [3.2]. Since the general effect of carbon is thus to increase strengthbut to decrease toughness, it is evident that the optimum carbon content must be chosenin terms of the properties required for a specific application.

3.6.2. Effect of Nickel

One of the major effects of nickel in the steels considered here is to increase their harden-ability. Another, and even more important one, is to lower the transition temperature sothat fracture at room temperature remains fully ductile, even at high strength levels [3.9];this effect, which must not be confused with an actual improvement of the impact strengthitself, is probably associated with the fact that nickel enhances cross-slip at high strainrates and/or low temperatures [3.19]. The effect of nickel in lowering the M? temperatureand thus in increasing the amount of retained austenite is counterbalanced by means ofcobalt additions.

35

COBALT-CONTAINING HIGH-STRENGTH STEELS

1.00 r

0.751

0.25

TEMPERING TEMPERATURE/C200 300 400 500

fig. 3.22. — Effect of silicon on toughness ofHP 9-4-45 in tempered martensitic condition.After J.S. P*si:ovm and S.J. MATAS J.?..'l.Condition : uustcnizatinn, oil quenching,refrigeration, and 2 -• 2 h tempering at indicated

temperatures.

300 400 500 600 700 800 900 ITOOTEMPERING TEMPERATURE, "r

3.6.3. Effect of Silicon and Manganese

The effect of silicon in reducing the toughness of HP 9-4-45 is illustrated in Figure 3.22.In addition to reducing the overall notch toughness, silicon shifts the 500 F (260 C)embrittlemeni to the 800 to 900 F (425 to 480cC) temperature range [.*..?]. Thus, specialmelting practices are used in order to keep the silicon content as low as possible.

Manganese is similar to nickel in so far as its effect on the transformation characteristicsis concerned. However, its presence in substantial amounts does not contribute to toughnessas nickel does indirectly. The upper limit for HP 9-4-25 is 0.35 %Mn, and most produc-tion heats did not exceed 0.30",; [3.20].

3.6.4. Effect of Carbide-Forming Elements

The basic role of the refractory-element additions during tempering of the 9Ni-4Co steelsis clearly revealed by comparing the curves in Figure 3.12 for the experimental 0.3C-9Ni-4Co and industrial HP 9-4-30 steels. The greatly improved tensile and yield strengths ofthe latter as compared with the former at all tempering temperatures are evident, as is thedecrease in impact strength measured after tempering. Other indications that carbide-forming elements reduce the toughness of the HP 9-4-45 steel and are conducive to SOOT(260'C) embrittlement have been provided [3.2]; this led to the recommendation thatthe chromium and molybdenum contents be kept to a minimum in HP 9-4-45 when thehighest level of toughness was required. These two elements do not have a detrimentaleffect on the toughness properties of the lower-carbon grades, HP 9-4-30 and HP 9-4-20,and provide enhanced strength, weldability, temper resistance and elevated-temperatureproperties.

The role of chromium and molybdenum has been investigated in greater detail in the caseof the IONi-8Co :,teel system [3.9]. Chromium was shown to shift the secondary hardeningpeak of the 0.12C-l0Ni-8Co-lMo steel to lower temperatures and to slightly higher values,while retarding softening. The optimum chromium content as regards secondary hardeningis 2%; higher levels result in too rapid coarsening of the M2C alloy carbides, while lowercontents are associated with problems of retained austenite, due to a higher secondary,hardening temperature. Molybdenum at the 1 % level is responsible for the occurrenceof a marked secondary hardening peak in the 0.12C-10Ni-8Co-2Cr steel; increasing themolybdenum content from 1 to 2 % does not significantly affect the tempering behaviour.Finally, vanadium at the low levels used in HP 9-4-X steels decreases the reaction ratesfor both the pearlite and bainite transformations; however, its major function is to act as agrain refiner [3.20].

36

3. CAKBIDE-SI RENGTHENED STEELS — PHYSICAL METALLURGY

3.6.5. Effect of Cohall

Some of the effects of cobalt in carbide-strengthened steels have already been mentionedin preceding sections. In brief, it ircreases the Mg temperature, refines the martensiticstructure, and leads to retention cf the dislocation substructure at higher temperingtemperatures, giving a finer precip lation of dislocation-nucleated alloy carbides. Theincidence of these structural features on the tempering response of cobalt-containing steelswill now be discussed in greater detail.

The effect of increasing cobalt contents on the yield and impact strengths of a 5Ni-Cr-Mosteel containing 0.25 %C was shown in Figure 3.13. Addition of 4%Co enhances thesecondary hardening effect, so that the yield strength is raised on tempering in the 800to 1000 F (425 to 540C) range [3.6]: increasing the cobalt content to 8",, increases theyield strength at all tempering temperatures. On the other hand, the notch toughness islowered, especially in the region of secondary hardening, so that the strength/toughnessrelationship for the Co-containing teels tempered at I000=F (540cC) is poorer than thatobtained for the Co-free one tempered at 400°F (205°C) [3.6, 3.7]. In the case of theformer, it has proved necessary to increase the nickel content and to adjust the carbon,chromium and molybdenum contents in order to achieve the optimum strength/toughnesscombination.

The effect of up to 8% cobalt on the tempering behaviour of a carbon-free, 10°oNi steelis to increase the hardness through a small solid-solution effect which appears to beretained over the whole tempering range [3.9]. Both the Co-free and the Co-containingsteels present a small yield-strength peak on tempering at about. 900 F (430;C).which is ascribed to the relief of residual stresses [3.21]. Cobalt additions to a0.12C-10Ni steel increase the hardness and strength values at almost all temperingtemperatures. In the presence of carbide-forming elements (Fig. 3.23), cobalt additions

TEMPERING TEMPERATURE,*C .WO 200 300 - 500 600

160 5 - 1

As quenched 200: UIO 600 ._ BOO1EMPERIN6 TEMPERATURE.°F

100040 3 .

1200

2QQO

180 £

160 <

U0E

120?

100?

0

SO 3

Fig. 3.23. — Effect of cobalt content on the tempering response of 0.12C-10Ni-2Cr-lMo steels. After L.F. PORTER et al. [3.6] and G.R. SPEICH et al. [3.9].Austenization : 15OO°F (815°C) - 1 h, W.Q. Tempering time : 5 hours.

37

i )H \1 I i O S I \ l \ l \ i i M l i i H M K I M J I I I S l L L U s

result m much higher a t tempered yield strength values; in particular, a sharper secondaryhardemns: peak is observed in the 0.1 2C-IONi-2Cr-l Mo sieei containing S",,Co than in'.he ctM\i!t-frec cme [.v''j. This etVect of cobalt in carbon-conlaming steels can be interpretedin terms oi the resulting liner dislocation structure (see Iig. 3.17).

A comparison between Figures 3.2.' and 3.13 reveals why better compromises can beachieved m the cohalt-modiiied lONi-Cr-Mo than in the eohah-moditied 5 \ i - C r - M osteels referred to earlier on: on lhe one hand, the increase in yield strength at the secondaryhardening peak is much more significant in I he former steels than in the latter; on theother, the toughness trough for the IONi-Co-Cr-Mo steels occurs ai a lower temperingtemperature than does the yield strength peak, with the result thai, at the temperaturecorresponding to the latter peak, the toughness has already risen sharply.

.viviv Srrcnatii Toughness vs. Sn'iiclure Ri'Luiunship

YicU •urennth. The v leld strength of low-carbon martensitic steels in the tempered conditioncan be considered as being derived from three primary strengthening mechanisms (seeChapie.' 2. Sections 2.3.6 and 2.-J.1) : (1) substruetural strengthening, resulting from thehigh dislocation density of the martensite. (2l carbide precipitation hardening, andI3I suh>tiiL,.ional-element solid-solution strengthening. In steels that are subject tosecondarv hardening, it is important to retain the dislocation substructure at temperaturesas high as possible since, in addition to its inherent strengthening effect, this structure willresult in a liner dispersion of dislocation-nucleated particles and hence in increasedprecipuat.on hardening. An estimation of ihe magnitude of the different strengtheningmechanisms operative in a 0.12C-10Ni-KCo-2Cr-IMo steel has shown that, for a yieldstrength of IS 5.000 psi (1200 MNm-'l. carbide strengthening contributes 45",1, sub-structural strengthening 37",,. and solid-solution strenthening 16";,, the balance beingaccounted for b> the fnclional stress of pure iron [3.9].

On the basis of the above considerations, it is possible to explain the role of the variouselements in determining the strength of tempered IO",,Ni martensites. Although alloyingis intended mainly to control the steel's tempering behaviour, the high nickel and cobaltcontents give rise to some solid-solution strengthening. Carbon obviously increases thestrength of the steel over the entire tempering range through an increase in the amountof carbides and the consequent decrease in intercarbide spacing. Chromium reduces iherate of softening but does not give a secondary-hardening peak in the absence of molyb-denum. The latter element (in association with chromium) is responsible for secondaryhardening through precipitation of MiC carbides.

The role of cobalt in increasing the temper resistance of alloy steels is well known [3.22].and has been dealt with in more general terms in Chapter 2 (Section 2.4.2). It can beattributed to two effects, vhich probably occur simultaneously ; (I) cobalt increases the

Y"DIFFUSION PATH FROM INTERFACE-^

oj partit ioning b) no partitioning

Fig. 3.24. — Interface mod>l showingthe effect of cobalt partitioningon carbide growth rate in steels.After V.K.. CHANDKOK et al. [3.22].Cci, is the concentration profilefor cobalt; or and ac(Co) are respectively the activities of carbon inthe absence and presence of cobalt.

38

i. CARBlbil-STRENGTHENED STEELS — PHYSICAL METALLURGY

dislocation density and retards the recovery of the dislocation substructure, so that thenumber of nuclcalion sites for the subsequent precipitation of carbides is increased [j.V];(2) it increases the activity of carbon, which results in a higher nucleation rate for. and aliner dispersion of, the carbide precipitates, as shown particularly clearly in the case ofhigh-speed steels [3.22, 3.23]. The elfeet of cobalt on the activity of carbon should alsolead to a reduction in the growth rale of the carbides, in conjunction with the partitioningof cobalt to the ferrite at the ferriie carbide interface: this effect is similar to th,.st of silicon.Figure 3.24 shows schematically how such partitioning leads to a lower ca:Son activitygradient at she interface (hence the decreased carbide growth rate) [.''.22].

Thus, cobalt enhances both substructure and carbide precipitation strengthening, i.e.. thetwo factors which have just been shown to account for the largest part of the strength inthe case of the 10Ni-8Co-Cr-Mo steel. However, the mechanisms mentioned in thepreceding paragraph as controlling the nucieaiion and growth of carbides do not ruleout a possible effect of cobalt on the diffusion of carbide-forming elements [3.24] or onthe surface energy of the carbide,ferrite interface [3.22]. The increased activity of carbondue to the presence of cobalt should also be responsible For the easier self-tempering ofcobalt-containing carbon martensites, and possibly for the increase in Mf temperature incarbon-depleted zones [3.22].

Toughness. As mentioned earlier, the role of nickel with respect to toughness is mainly todecrease the transition temperature, and hence to ensure retention of high room-temperature toughness at high strength levels; there is. however, a limitation to the useof nickel as an alloying element, which is primarily related to its tendency to increaseaustenite retention or reversion. The effect of the other alloying elements on toughnesscan be interpreted in terms of their role in promoting the substitution of the much tinerMiC for MjC: ductile fracture occurs by nucleation of voids at particles, and suchnucleation is more difficult when the par ides are small [3.9\.

3.7. Concluding Remarks

The development of high-alloy steels with yield strengths typically in excess of 140,000 psi(1000 M N/m2) has been characterized by the consistent efforts made to attain simultaneouslyhigh toughness values. In addition to a general optimization of carbon-containing high-strength steels such as the AISI 43xx series and H-l 1, this aim has been pursued essentiallyalong two lines : (1) the development of low-carbon martensites strengthened by inter-metallic compounds, and (2) the development of steels with the highest carbon contentscompatible with high toughness and good weldability.

The present chapter has been concerned with the second type of development whichcharacteristically involves the use of high nickel additions to achieve high toughness.Imperatives regarding the steels" transformation structures {e.g. restriction of the retainedaustenite content) led to associating this element with cobalt: the levels which were selectedon the basis of property requirements, and possibly cost, were 4 and 8"oCo for theHP 9-4-X (9%Ni) and lONi-Co-Cr-Mo steels, respectively.

The desire to ensure good resistance to tempering led to the use of the carbide-formingelements chromium and molybdenum. In the lONi-Co-Cr-Mo steel, the optimum chromiumand molybdenum contents were determined as 2 and 1 %. respectively, whereas in theHP 9-4-X steels, they were held to 1 % or less depending on the carbon content. Thelatter steels also contain a small vanadium addition (0.1 %). Cobalt appears to be a useful,if not essential, element in controlling the tempering behaviour.

As regards the carbon content, the level chosen for the lONi-Co-Cr-Mo steels is ratherlow (0.12%), siuce this type of steel was initially intended for heavy-section welded

39

lOBAl r-l O \ r USING IIK.II-sriU-NGTII STKtLS

>irucuires exhibiting high toughness (room-temperature Charpy Y-notch \ahies greaterthan wH'i-llv. j.t . . MU) at \lcki strength le\els up to 21)0,000 psi (1400 MN..-m:). Thecarbon level of H P l M - \ steeK varies from 0.20 to 0.45",, depending on the combinationof toughness, weldability and strength that is required.

The aood properties of these steels in the tempered marten-.itc condition result from thetact that their >.iruciure> are strengthened b\ a line dispersion of alloy carbides whichsubstitute at least partially for the embrittling cementite precipitate. This process isassociated uub a relali\e inseii>itivitv of pri.perties to tempering temperature in ihe caseof the HIJ lJ~i-\ steels and. in that of the !U\i-(.'o-C"r-Mo steels, with the existence of asecondary ha; Jening range uver which both the yield strength and toughness increase.The former sleek, particularly vhe HP"-M-45 grade, also exhibit excellent combinationsof toughness and strength in the lower-bainitic condition.

4. CARBIDE-STRENGTHENED STEELS — PROCESSING AND PROPERTIES

This chapter deals mainly with the processing, properties and use pattern of the HP 9-4-Xsteels. Data on the lONi-SCo-Cr-Mo steel will also be included, although this steel is notso well documented, due partly to the restricted distribution of some of the reports on itidevelopment. However, in view of the similarity between the IONi-8Co-Cr-Mo steel andthe lower-carbon HP 9-4-X grades as regards their physical metallurgy (see Chapter 3).an attempt has been made to give an overall picture of the former, in spite of the absenceof specific data in many cases.

4.1. Primary Processing

The HP 9-4-X steels are produced by electric-furnace air melting pins consumable-electrodevacuum-arc remelling. The optimum combinations of properties are obtained using theVAR-CDOX process (vacuum-arc remeiting plus carbon deoxidation): other meltingpractices (air melting plus Si Al deoxidation: air melting, Si/Al deoxidation and vacuum-arc remeiting) can be used when the requirements are less stringent [4.1]. Specifically,the VAR-CDOX process involves electric-arc melting an unkilled heat having excesscarbon, followed by consumable-electrode vacuum-arc remeiting using a cold-walledcrucible [4.2]. Carbon deoxidation occurs during the second step, i.e., under vacuum:this favours elimination of the gaseous reaction product. It has in fact been shownthat steels prepared in this way tend to be cleaner than those that are air melted, or airmelted plus vacuum remclted. and to have lower gas contents. Both factors lead to animprovement in the toughness of the steels [4.3].

As regards the lONi-SCo-Cr-Mo steel, the melting practice reported consists of vacuum-induction melting followed by vacuum-arc remelting. and involves vacuum carbondeoxidalion [4.4. 4.5].

The hot-working practices used for HP 9-4-X steels are similar to those for AISI 4340[4.1]. The maximum heating temperature has been fixed at 2O5OLF (1120'C) for carbon-deoxidized material, as rapid grain coarsening can occur above this temperature in theabsence of significant amounts of aluminium: in Si/AI deoxidized material, graincoarsening has not been observed below 2150 F (1175 C) [4.6]. When forging VAR-CDOXingots of HP 9-4-45 and HP 9-4-25, a maximum reduction of 75% and a maximum finishtemperature of 1900-1950F (I040-1065C) have been recommended.

The stand.—,! heat treatments for the HP 9-4-X steels are listed in Table 4.1. As regardshardenability, essentially no drop in hardness with distance from the quenched end ofJominy bars has been observed for either the 0.25C or the 0.45C grade; this indicates

40


TAB.' ' •.:. - STANDARD HEAT TREATMENTS FOR HP9-4-X AND lONi-SCo-Cr-Mo STEELS

HP 9-4-45 [4.7]Normalize al IWI0-I65O I- (S70-900 Cl for I hour per inch (min. I hour), air cool.Ausleni/e al 1475 : 25'F (790-815 C) for I hour per inch (min. I hour).

Martensite Hamilt-on quench refrigeration at —100 F ( — 73 Ci for 2 to ft hours. Quench in salt at 4f>5 : 5 F (240-245 Ci . hold atTemper at 400 F (205 C) for 2 • 2 hours. temperature for 4 to X hours, air cool.

HP 9-4-30 [4.7]i. Normalize at I (-50-1700 F (900-930 Cl for I horn per inch (min. I hour), air cool.

Austeni/e at 1550 : 20 F (830-855 Cj for I hour per inch (min. I hour),Marwrtsilt' tiuinile

Oil or water quench • refrigeration al 100 F ( 73 Cl for I to 2 hours. Quench in salt at 450 5 F (230-235 Cl . hold ali Temper a! 1000 F (540 C) for 2 ; 2 hours. temperature for 6 to 8 hours, air cool.

HP 9-4-25 [4.7],: Normalize at 1650-1700.-F (900-930 C) for 1 hour per inch (min. I hour), aif cool.! Auslenize at 1550 .•;. 20 F (830-855 C) for 1 hour per inch (min. I hour)., Oil or water quench.

l.i Temper al 1000 F (540 C) for 2 : 2 hours.

HP 9-4-20 [4.8];i Normalize at 1650 ± 25 F (885-915'C) for 1 hour per inch (min. 1 hour), air cool.•i • Austenize at 1500 i 25' F (800-830 C) for 1 hour per inch, water quench..•; Temper at 1025 •;• 25:F (54O-5651 C) for 4 to 8 hours (min. 4 hours).

10Ni-8Co-Cr-Mo [4.9]11 Normalize at 1660F (905 C) for 1.5 hours, water quench.'i. Auslenize al 1500"F (815'C) for 1.5 hours, water quench.

i Temper at 950 F (510 C) for 5 hours.

that sections at least 3 in. (7.5 cm) thick can be fully hardened [4.10]. The hardenabilityof HP 9-4-20 and HP 9-4-30 appears to be even larger : according to the CCT curvesgiven in Chapter 3 (see Section 3.2.1), the rounds that can be fully hardened by waterquenching should reach 4 and 12 in. (10 and 30 cm) in diameter, respectively.

The strongest grade in this family, HP 9-4-45, can be strengthened by a martensitictreatment (the A/s temperature is about 450°F, i.e., 230°C), but this must include arefrigeration step to minimize the retained austenite content. It can also be subjected to abainitic treatment, which produces higher toughness but a somewhat lower yield strength.The bainitic transformation in the 450-650T (230-345°C) temperature range is completedin less than 2 hours; the 4 to 8 hours mentioned in Table 4.1 are primarily a precautionarymeasure, particularly in the case of large sections [4.10].

The HP 9-4-30 steel may be quenched to martensite or isothermally transformed to bainite:however, the martensitic heat treatment is normally recommended for this grade, since thebenefits of the bainitic treatment are considered as marginal.

The HP 9-4-25 steel, which was developed primarily as a weldable alloy, has Ms and Mf

temperatures of approximately 590 and 480°F (310 and 250'C); these relatively hightransformation temperatures allow self-tempering of the mariensite to occur, leadingto good toughness in as-quenched or as-welded material. It is not possible to produce afully bainilic structure in this grade over a wide range of reaction temperatures as was thecase for the higher-carbon steels; accordingly, the only heat treatments used are thosethat will produce a martensitic structure. However, such treatments are adequate sincethe toughness of the tempered martensite is sufficiently high and would not be significantlyimproved through a bainitic treatment.

Similarly, the lowest-carbon grade, HP 9-4-20, is always used in the martensitic condition;the double or single temper at 1025°F (550°C) provides an excellent combination of

41

LOHAl . I -LOMAIMMJ HK.H-SIKBNUIH Slfcfc'LS

strength and fracture toughness. This steel, which was developed for maximum weldability,is also amenable to self tempering.

As regards the lONi-SCo-Cr-Mo steel, it is normally used in the quenched and temperedcondition. The standard heat treatment listed in Table 4.1 involves a double austenizingtreatment [4M], though this has occasionally been replaced by a single treatment at 1550or 150H F tS4? or i>15 C) [4.11, 4.12}. Hardenabilitv is more than adequate for platethicknesses up to at least 4 in. (10 cm) [4.4\.

4.2. Properties

4.2.S. Strength Toughness Cluiructcristics

Representati\e properties of the steels considered in this chapter, in their standard heat-ireated conditions, are listed in Table 4.2. It is seen that the highest strengths are obtainedin the HP 9-4-45 and HP 9-4-30 grades and are associated with good toughnesses, parti-cularly in the hainitic condition; for instance, A,'u. values of 65,000 to 95,000 psiy in.t?l to 104 MNm -; -) and IIV.-I precracked Charpy impact values of 1000 in.lb/in.-(176 kJm-l have been reported for forged HP 9-4-45 (bainitic) at an ultimate tensilestrength level of 270.000 psi (1860 MN/m-) [4./]. The lower-carbon grades are conspicuousfoi their very high toughness values at reasonably high strength levels. For instance,HP 9-4-20 at the 200.000 psi (I3S0 MN/m-) ultimate tensile strength level, shows a typicalA',,, value of 155.000 psi\ in. (170 MNm-W), while values of up to 183,000 psiy in.i.2t)l MNm ; -) have been reported [4.8].

On the basis of Charpy V-notch data, the IONi-8Co-Cr-Me steel exhibit? •' better combination of strength and fracture toughness than does HP 9-4-20 [4.9]. K]c values determinedaccording to ASTM specifications are not available for this steel, because of the largespecimen thickness required. Estimations based on results of tests on sub-thicknessspecimens are in excess of 250.000 psiy in. (275 MNm-*'-) [4.13], and Ku. values as highas 300.000 psi\ in. (330 MNm ••<•'-) have been cited [4.5].

The effect of melting and clenx'ulaiiun practice on tbe strength/toughness relationship hasbeen investigated in the case of the HP 9-4-45 steel. Figure 4.1 clearly shows the improvedtoughness of vacuum carbon-deoxidized heats as compared with air-melted Si/Aldeoxidized ones.

As regards the effect of heat treatment. Figure 4.2 shows the influence of temperingtemperature on the properties of the HP 9-4-20 and HP 9-4-30 grades. The strength/

TABLE 4.2. — TYPICAL ROOM-TEMPERATURE MECHANICAL PROPERTIES OF HP 9-4-X AND 10Ni-8Co-Cr-Mo STEELS

Materialand condition *

HP9-4-45 (muncnsitic)HP9-4-45 (bainuiclHP9-4-3O(martsnsilic)HP9-4-30 (bainuicj

HP 9-4-25HP 9-4-20

lONi-SCo-Cr-Mo

U

10-' psi

2X0-300260-280220-240220-240195-210190-215

196

T.S..

MS m~

1930-20701790- IV301520-16601520-16601345-14501310-1480

1350

0.2'

HP psi

245-260220-235190-200190-200178-192180-195

185

„ Y.S..

MWm*

16F.O-I7901520-16201310-13801310-13801225-13251240-1345

1275

Elong.in 2 in..

6-1012-1412-1612-1615-Io14-20

22

R.in A..

20-3540-5050-6050-6055-6561-71

71

C.V.N

ft.lb

10-1220-3020-2525-3032-4051-72

81

impact,

J

14-1627-4127-3434-4143-5468-98

110

Kir,

SO psi v in- ^

50-7065-95

100-120120-135> 140

155-175300**

55-77

71-104110-132132-141> 154

170-192330**

Hard-ness,Re

51514444424245

Ref.

[4.7]

[4.7]

[4.7]

[4.1,4-7]

[4.7]

[4.7,4.8]

• [4.5]

* Heat treatments us in Table 4.1." Value computed from results obtained on nan-siandaril specimens.

42


YIELD 5TR6NGTH: MN/m>I 5 o o ; •- •••- -is

TEMPERING TEMPERATURES

910 . 220 r -230 , • ;2A0YIELD STRENGTH, tfpsr:'

Kig. 4.1. — EITecl of deoxidation practice on thestrength/toughness relationship in HP 9-4-45 sheet,l-atigue cracked centre-notched specimens 0.180 in.!4.h mm) thick were used in determining the longit-udinal notched strength. After SJ. MATAS [4.3].

- lam

As « 400quenched

800 1200 » 400TEMPERING TEMPERATURE.T

800 1200

a) HP 9-4-20. Single 4-hour temper. l» HP 9-4-30 (martensitic).Double 2^2 hour temper.

Fig. 4.2. — Effect of tempering temperature on properties of HP 9-4-20 andHP 9-4-30. After R.T. AULT [4.14] and J.S. PASCOVER [4.15], respectively.

toughness relationship for the HP 9-4-45 steel in both the martensitic arid bainiticconditions has been plotted in Figure 4.3; the superior toughness of the bainitic structure,at a given strength level, is evident; this is maintained down to fairly low test temperatures,as will be shown further on (see Fig. 4.9). The effect of tempering temperature on theyield strength and toughness of the 10Ni-8Co-Cr-Mo steel was described in Chapter 3(see Fig. 3.23, p. 37).

Finally, the possibility of improving the stress/toughness relationship through thermo-mechanical treatment has been investigated for the HP 9-4-45 (martensitic) and HP 9-4-25grades [4.17]; the effect of strain tempering was evaluated in both cases in terms of theextent of cold work (roiling) and the tempering temperatures prior to and after cold

Fig. 4.3. — Relation between strengthand impact toughness of 0.5 in. HP 9-4-45 plate heat-treated to bainitic andtempered martensitic structures. AfterJ.S. PASCOVER and S.J. MATAS [4.16].Data points from left to right corres-pond to decreasing bainitic trans-formation temperatures or decreasingmartensite tempering temperatures.

YIELD •STRENGTH, MN/m '

1000 1200 1400 1600

KO 160 180 MO 220 240 260YIELD STRENGTH, B 3 psi

43

Kil l \LT-( MNTAIMNti HKiH-M IU-MI 1 H STl'.l-lS

PRETEMPERING TEMPERATURE. 'C....200'.. X » -400^ . 500

RETEMPER1NG TEMPERATURE/C300 j}- 400 : ; : 5 0 0 6CXT

0 n 20 30THICHNESS REDUCTION, V.

40 300 500 700 900 1100. PRETEMPERING TEMPERATURE. "F

'RT 200 " « 0 600 . SOO WX) 1200RETEMPERING TEMPERATURE,-F .

EtTccl i)l' .mioum of deformation Ipre- »*>) Effect of pretcmperinK ;cmpc: .i.'-irc i -5 ' \ ,d posl-U'mrcrmg ai 4\XT F. i.e. 205 Ci. deformation, no posl-deformult'Jit ijinpcring}.

ri Effect of retenperiny tcmpemturL' (pre-icmpuringat 30'J' F.i.e., 15O5C.25"„deformation).

hii; -i.i. — Effect of strain tempering on mechanical properties of HP 9-4-45 in the martenshie condition. After D.-KAUSII el a!. \4.17\

working. The results obtained for HP 9-4-45 are presented in Figure 4.4. Cold work wasfound to produce a steep increase in yield strength within the first 5",, of deformation,subsequent straining giving rise to a more moderate increase. Pretempering at lowtemperature (around 300 F. i.e.. 150 Cl resulted in much more pronounced strainhardening than did pretempering at high temperature (about 1100'F, i.e., 595QC); this canbe explained in terms of the carbide morphology developed during the pretemper, viz.finely dispersed =: carbides at 300" F and coarse cemerttite precipitates at 1100cF. Retemper-ing after cold working gave rise to a 20.000 to 30,000 psi (140 to 205 MN/in2) increase inyield strength, independently of the degree of deformation and the tempering temperatureover the 200 to 400 F (95 to 2O5X) range. All these effects are strikingly summarized inFigure 4.5a which shows the KUJY.S. relationships for the steel both after the conventionalmartensitic heat treatment and after subsequent strain tempering; although the two curvesinitially coincide, that for the strain-tempered condition extends up to very high strengths(of the order oi 380.000 psi, i.e., 2600 MN/m2), which are associated with acceptableFracture toughness (around 50,000 psi\ in., i.e. 55 MNm"3 '2). HP 9-4-25 can also bestrain tempered to extremely high yield strength levels, and these prevail in conjunctionwith higher toughnesses than are obtainable with HP 9-4-45 (Fig. 4.5b).

YIELD STRENGTH MN/m*_2«0 Z300 1S00

42D 180 220YIELD STRENGTH, 103psi

Fig. 4.S. — Strength/toughness relationship for HP 9-4-45 (martensitic) and HP 9-4-25 in boththe conventionally heat-treated and strain-tempered conditions. After D. KALISH el al. [4.17].

4. CARBIDE-STRENGTHENED STEELS - PROCESSING AND PROPERTIES

ID 20 30 40 50THICKNESS REDUCTION, %

Fig. 4.6. — Effect of strain tempering on mech-anical properties of HP 9-4-45 in the bainitiecondition. After D. KALISH el al. [4.18].

Transformation to bainite at 500'FCttfTC); pretempering at 400°F (2O5=C);

no post-deformation tempering.

Fig J.7. — Comparison of strength/toughness relation-ships for the HP9-4-X and other high-strength steels.Alter [4.8] for HP 9-4-20, J.S. PASCOVER [4.15] for HP9-4-30, J.H. GROSS [4.13] for I0Ni-8Co-Cr-Mo, and J.S.PASCOVER and S.J. MATAS [4-16] for the other steels.

260

240

220

200

ISO

. 160

^•140sa

120

100

80

SO

40

20

1200YIELD STRENGTH MN/m*

1400 1600 1600T

10Ni-8Co-Cr-Mo

2000 2200

. 18Ni-Co-MoMARAG1NG STEELS ;

250

200

I50 :

100

50

ISO 180 ZOO 220 240 260 280 300YIELD STRENGTH. 103psi

The effect of strain tempering HP 9-4-45 in the bainitic condition has also been investigated[4.18~\. In the same way as for the martensitic condition, strain tempering lower bainiteleads to increased strength and decreased toughness (Fig. 4.6).

Ausforming HP 9-4-45 has also been shown to increase the strength properties and todecrease toughness {4.17}. However, due to the fibering effect associated with rolling inthe austenitic condition, anisotropy in fracture toughness develops; the resulting longi-tudinal strength/toughness curve is close to that for the steel in the conventional andstrain-tempered martensitic conditions, but the transverse strength/toughness data pointslie below it.

This section can be usefully concluded by comparing the strength/toughness relationshipsfor the HP 9-4-X steels with those of other high-strength steels (Fig. 4.7). It is apparentt!i*.t the toughness of the KP 9-4-X steels is significantly greater than that of other carbide-strengthened steels such as H-l 1 or AISI 4340. At a strength level of 190,000 psi(1310 MN/m2), HP 9-4-20 exhibits Klc values of up to 180,000 psi\/in. (198 MNm"3 '2),while HP 9-4-30 appears to be at least as good as the 18%Ni maraging steels at strengthlevels of "200,000-210,000 psi (1380-1450 MN/m2). Although comparable K]c data are notavailable for the 10Ni-8Co-Cr-Mo steel, the values quoted in Table 4.2 indicate that, atthe 190,000 psi (1310 MN/m2) level, the curve for this steel probably lies well above thehighest values for HP 9-4-20.

45

l O H M . r - t O N T A l M M I HK.H-S I Rl-NCH H STEKLS

TE5T- TEMPERATURE.'C,0_ 200 400 600 800 BOO

JO

BO

60 j

40

20 =—-cSS!!gil>'n

r '•Q

/

No'

0 400 800 E00 -BOO 2000TEST TEMPERATURE. - F

TEST TEMPERAHJRE,"C.-200 " -SO • -100 -50- 0

TEST TEMPERATURE,*C-200 -100 0

-»a -2oo -wo otEST 1EMPERA1URE/F

..-300 -200 -100.. 0 BOTEST TEMPERATURE, T

Fig. 4.8. — Temperature dependenceof tensile properties of HP 9-4-20.

After R.T. A L L T [4.14].Condition: P 0 0 F ( 9 3 O C ) - l h . A . C . - 1500 F<8I5C) - th. W . Q . - IO5OF (565;C) - 2h. A.C.

ii) Tensile properties and frac-ture toughness of HP 9-4-20plate. After C. VISHMEVSKY..id E.A. SrEictRWALO y.iV].Condition : standard ht. treat.

h) Impact strength of HP 9-4-45 and HP 9-4-20. Respectivelyafter J.S. PASCOVER and S.J. MATAS [4.16]. and R.T. AULT [4.14],

HPy-4-45 was treated to 250.000 psi (1720 MN/m')U.T.S. in both the bainilie and martensitic conditions:HP 9-4-20 received the conveniional heat treatment.

Fig. 4.9. — Properties of HP 9-4-20 and HP 9-4-45 at low temperatures.

4.2.2. High- and Low-Temperature Properties

Since the steels discussed in this chapter are essentially intended for use in the neighbour-hood of room temperature, information regarding their high-temperature properties israther scarce. Figure 4.8 illustrates the effect of test temperature on the mechanicalproperties of HP 9-4-20 after its standard heat treatment.

Regarding the properties of '>ese steels at cryogenic temperatures, F'gure 4.9a showsthose determined on HP 9-4-20 plate in the heat-treated condition. The ultimate tensilestrength is seen to increase from 200,000 psi (1380 MN/'m2) at room temperature to

- 2 0 C

1000

Fig. 4.10. — Stress-rupture and 0.2",, plas-tic deformation creep curves for forgedHP 9-4-20 bar (transverse). After [4.20].

10* TO5 10s

CYCLES TO FAILUHE

Fig. 4.11. — Smooth and notched fatigue behaviourof HP 9-4-20 (Moore rotating-beam fatigue tests ontransverse specimens). After R.T. AULT [4.14].Specimens tempered at IOOO1VF (54OSC) lo ultimatetensile and yield strength levels of 210.500 and187.000 psi f 1450 and 1290 MN/m*), respectively.

46

4 < \ R H I D I - S I K I - N ( , ! I I I - N [ . D M l I I S P R O C E S S I N G A N D P K O P F . R 7 l i : s

Fig. 4.12. — Dependence of fatigue crackgrowth rale in air on stress intens-ity factor range lor HP 9-4-20 andlONi-XCo-Cr-Mo steels. After [-4M].

o HI'9-4-20; lONi-KCo-Cr-Mii

STRESS INTENSITY FACTORRANGE^K.MNmWjjj>. '" ;-- ' '• '•• '• ' 5 0 '1-,-X' ' 1 0 0 :-•-...• 2 0 0 . A > - i ' . i •

£ 2020 50 100 200STRESS INTENSITY FACTOR RANGE, &K.103 psi \fin.

255,000 psi (I760MN/m-) at —32OCF (—I96°C), at the expense of a sharp decreasein fracture toughness and a slight decrease in ductility [4.19]. Similarly, the ultimateteusile strength of the IONi-8Co-Cr-Mo steel increases from 204.000 psi (1400 MN/m-)to 266,000 psi (1830 MN/m2) over the same temperature range [4.12]. The effeci oftemperature from —300 to 120F (—185 to 50;C) on the impact strength of HP 9-4-45and HP 9-4-20 is shown in Figure 4.9b. As was to be expected from the typical behaviourof high-nickel steels, the decrease in impact strength with decreasing temperaiures isnot characterized by a transition temperature.

Finally, creep and stress-rupture data determined on forged HP 9-4-20 bars are presentedin Figure 4.10.

4.2.3. Fatigue Behaviour

Results of room-temperature axial fatigue tests on both smooth and notched specimensof HP 9-4-20 are plotted in Figure 4.11. Further data on the fatigue behaviour of HP 9-4-Xsteels will be found in the literature [4.5, 4.10, 4.20]. The fatigue crack propagation rateof several high-strength steels tested at room temperature in an air environment wasshown to obey the following relationship : do/dW = 0.66 x IO~8 (A/C)2-25 [4.21]. Figure4.12 illustrates this relationship in the case of HP 9-4-20 and 10Ni-8Co-Cr-Mo, thoughin this case the exponent of AK was found to be 2.5.

4.2.4. Stress-Corrosion Characteristics

As is the case for other high-strength steels, the stress-corrosion resistance of the HP 9-4-Xsteels determined on notched specimens is generally improved by factors that tend toincrease its fracture toughness, such as reduced strength level and carbon content, as wellas by special melting practices which minimize non-metallic inclusions [4.22].

Threshold intensity factors KUcc of several high-strength steels in 3.5 % sodium chloridesolution are given in Table 4.3. It is apparent from these results that HP 9-4-45 has onlymarginal resistance to stress-corrosion cracking in the martensitic condition; it is definitelymore resistant in the bainitic condition, though still greatly inferior to HP 9-4-30. Theseobservations confirm the above-mentioned relationship between toughness and stress-corrosion resistance. The KUcc values for the steels HP 9-4-20 and 10Ni-8Co-Cr-Mo

47

I OHM 1 l ONI MMMi HKiH SIRFNl.TH STEELS

190]

10010 3To 400 600 800 BOO 1200 K00 1600 1800

TIME 10 FAILURE, hours

Fit . 4.13. — K\,rr stress-corrosionbehaviour of HP 'MOO in artificia'sea-water (3.5",,NaCl). After y.S\.

r> Air tesis• Sea-water tests• > No failure after six months

reflect the latter's superior stress-corrosion resistance. Results of stress-corrosion testson HP 9-1-20 in synthetic sea water are shown in Figure 4.13.

4.3. Secondary Processing

The lower-carbon HP 9-4-X steels are generally formed after heat treatment [4.25]; thisalso applies to the HP 9-4-30 grade treated to martensite, but not to HP 9-4-45, which wasfabricated in the annealed condition. Possible operations include rolling, drawing, bending,roll forming, shear spinning, explosive forming, etc. The effect of cold working on themechanical properties of these steels was described in Section 4.2.1.

All four HP 9-4-X steels are easily machined in the annealed condition. After heattreatment, their machinability is comparable to that of AISI 4340 at similar hardnesslevels [4.7].

Finally, the HP 9-4-X steels, in particular the lower-carbon grades, have been shown topossess excellent weldability. Data have been reported for gas tungsten arc (TIG), hot-wireTIG, alternating-current gas-metal arc (MIG) and electron-beam welding procedures[4.1]. The lower-carbon grades can be welded in the heat-treated condition, using fillermetals which closely match the base metal composition [4.26]; tbey require ho post-weldheat treatment in view of the good toughness of the self-tempered martensite and/or

TABLE 4.3. — THRESHOLD STRESS-INTENSITY FACTOR IN 3.5% SODIUMCHLORIDE SOLUTION FOR DIFFERENT HIGH-STRFNGTH STEELS

11

II

Steel

300 M

HP9-4-45 (martensiticl

HP 9-4-45 (bainitie)

4330 V

H 11

HP 9-4-30 (martensitic)

18Ni(250) maraging

HP 9-4-20

lONi-SCo-Cr-Mo (plate)

l0Ni-8Co-Cr-Mo (hoi pressed)

10^ p .

283

276

266

239

219

231

259

205

196

203

U.T.S.,

i MNjiri1

1951

1903

1834

1648

1510

1593

1786

1412

1350

1400

y.s..103 psi i

236

236

220

196

188

200

249

190

185

191

1627

1627

1517

1351

1296

1379

1717

1310

1275

1317

103psiv'in.

76

69

89

103

54

116

92175

.100

> 234

f

S3

76

98

113

59

127

101

192

330

>257

103psi\/in. M

13

15

20(est.)

25

30

45 (est.)

45

110

180210

14

16

22

27

33

49

49121

198

231

Reference

[4.23]»

»

»

»

»

»

[4.8]

[4.5]

14.24]

48

4 C ARBiDI-.-STR!-N<,rHr.N[-:n S'IkF.LS f'ROC LSSINd AND PROPERTIES

athermal bainite which form on cooling [4.27], In the case of the HP 9-4-25 grade, jointefficiencies of 95 to 100",, have been reported [4.1). The properties of HP 9-4-20 we-dsarc equivalent to those of the heat-treated base plate [4.14]. In order to overcome detri-mental effects of stressrelieving treatments on the toughness of the welds, studies on theuse of modified filler-metal compositions were recently carried out; they resulted in thedevelopment, in the case of HP 9-4-20,of a completely stress-relievable system, i.e., one inwhich the base metal, heat-affected zone and weld are totally free from stress-reliefembrittlement [4.28]. As regards the HP 9-4-45 grade, it was shown that, like othermedium-carbon ultrahigh-strength steels, it must be welded in the annealed or normalizedcondition and then fully heal treated to develop the desired properties; using filler metalsof matching composition except for a slightly lower carbon content, joint efficiencies ofnearly 100",, on the basis of yield and ultimate tensile strengths were achieved, withfracture toughnesses generally greater than 80",, of that of the base metal [4.29].

The formability and machinability of the IONi-8Co-Cr-Mo steel should be comparableto those oi' the lower-carbon HP 9-4-X grades. In addition, the steel has been successfullywelded by the TIG process using weld wire of composition nearly matching that of thebase metal [4.9]. A comparison of weldments made v/ith iuNi-8Co-Cr-Mo and HP 9-4-20wires showed that the most desirable weld properties are obtained using the matchingweid wire and a tempering cycle. Under these conditions, tensile properties, impactstrength, precracked impact energy and " apparent " fracture toughness of the weld centreand heat-affected zone are close to those of the base metal. It is stated, however, that theimprovement gained through retempering is probably not sufficient to justify the additionalcost in most cases.

4.4. Applications

The applications of the HP 9-4-X steels were reviewed fairly recently [4.25]. The higher-carbon grades are usually preferred for components which do not require welding duringfabrication or which do not preclude a final heat treatment; this is particularly so forsmall parts which are fabricatedby forging, machining, and possibly welding prior toheat treatment. For instance, the remarkable combination of strength, toughness andfatigue resistance of HP 9-4-45 made it an ideal choice for forged parts in critical aero-nautical applications. The steel was used for fasteners and aircraft landing-gear links,where resistance to fatigue is an essential requirement; connecting rods and valve springwire for racing cars were also made from this grade. HP 9-4-30 [4.25] and 10Ni-8Co-Cr-Mo[4.30] have been used in light-weight dual-hardness armour and heavy-section homogeneousarmour plate because of their good hardness/toughness relationship. On account of itshigh structural stability up to at least 800°F (425°C), associated with satisfactory retentionof mechanical properties, HP 9-4-30 has also been selected for several forged structuralcomponents on the Boeing 747 aircraft [4.25). HP 9-4-20 has also found applications asforged components in advanced aircraft, due to its high fracture toughness [431].

For applications which require welding during fabrication, the HP 9-4-20 and 10Ni-8Co-Cr-Mo steels have the advantage over other commercial alloys at this strength level thatthey can be welded in the heat-treated condition and require no post-weld treatment.Moreover, their deep-hardening characteristics, together with their high strength, highcrack-propagation resistance (ATlc. and îjCC) and superior toughness, make them idealfor use in heavy sections in aircraft, aerospace and hydrospace applications. In fact, both10Ni-8Co-Cr-Mo and HP 9-4-20 have interesting potentialities for the manufacture ofinternally or externally pressurized vessels, such as reactors and submersibles. In thisrespect, a great deal of evaluation work has been devoted to HP 9-4-20 by the U.S. AtomicEnergy Commission and the American Society of Mechanical Engineers in connectionwith nuclear reactors and pressure vessels, respectively.

49

1 - i O M A I M M I H I C H S I RFNCi lH S T E M S

? Ni-Co-Mo MARAGING STEELS — PHYSICAL METALLURGY

5.1 Background

The maraging steel:- are a relatively new class of ultra-high-strength sieels that dcrisetheir strength from hardening mechanisms other than the classical ones associated withcarbon martensite. bainite or precipitation of carbides during tempering. These steels.which possess combinations of strength and toughness that are among the highestattainable in commercial alloys, are characterized by their very low carbon content andihe use of suhsiituiionul elements to produce age-hardening in Fe-Ni mariensites.

The metallurgical principle on which manging steels are based was established as earlyas 1939 [5.1] : the thermal hysteresis between martensite formation in Fe-Ni alloys oncooling and its reversion to austenite on heating, and ihe increase in this hysteresis withincreasing Ni content (Fig. 5,1). It can be seen that the reversion temperature of a 20"v,Nialloy is about 1IO5;F (595 C). which is sufficient to allow ageing of the martensitic matrixat about 900 F (480C). The term " maraging " was coined to indicate that theprecipitation reactions that are responsible for the ultra-high strength of these steels occuron ageing them in the martensitic condition.

The actual development of rnaraging steels was carried out at the Internationa! Nickel Co.in the late I950's [5.3]. The early work ied to the first two grades of maraging steel, theso-called 20",, and 25"0Ni steels [5.4). The A/., temperature of the alloys was controlledby adjusting the nickel content. After formation of the low-carbon martensitic structure,these steels, which contain a combination of 0.3%AI, I.4".,Ti and 0.4°uNb, wereprecipitation hardened during ageing between 800 and 950F (425 and 510 C). Hardnessesas high as Rc 67 were obtained, and good combinations of strength and ductility athardness levels of Rc 53-56 were reported.

However, the 20 and 25",,Ni grades were soon abandoned on account of their brittlenessat ukra-high-strength levels [5.5]. They were replaced by steels of the l8Ni-Co-Mo type,which have successfully withstood the test of time. The appreciable hardening of Fe-Ni

Fig. 5.1. — Fe-Ni transformation diagram.After F.W. JONES and W.I. PUMPHREY [5.2],

50

5. Ni-Cu-Mo MAKAGINd SIKhl.S PHYSICAL METALLURGY

20 30 40wt.%Coxwt.%Mo

60

Fig. 5.2. — Effect of cobalt • molybdenumproduct on maximum hardness of Fe-18.5 to20.1 '•„ Ni alloys. After R.K. D K KF.R el at. [5.6).o annealed I h al 1600 T (870"Cl, A.C.* maraged 3 to 10 li at 800-900 F (425-4hl0C).

12001 (2501 13001 13501 i «B I 15001MARAGING STEEL GRADE

Fig. 5.3. — Ultimate tensile strength,fracture toughness and cobalt andtitanium contents vs. maraging steelgrade. After A. MAGNEE er al. [5.7].

martensites that occurs when combined additions of cobalt and moKlxlenum are madewas reported in 1960 [J.6]. Figure 5.2 illustrates the synergistic effect of coball withmolybdenum on the age-hardenability of Fe-18"0Ni alloys. This binary composition waschosen as the alloy base since higher nickel contents led to retained austeniie. A numberof l8Ni-Co-Mo alloys were developed in rapid succession; later on. other Ni-Co-Mocompositions with less nickel were introduced.

Basically, the heat treatment of maraging steels consists in solution annealing for I hourat 1500°F (815"C), although other annealing temperatures or multiple annealing treatmentsare used in certain cases. Upon cooling in air to room temperature, the alloys transformcompletely to martensite. Because of their high nickel content and the virtual absenceof carbon, hardenability is not a problem and the cooiing rate after annealing is unim-portant. In the as-annealed condition, the alloys have a hardness of the order of Rc 30.and can be readily machined or fabricated. Hardening is then achieved by maraging,generally for 3 or 6 hours ;tt 900cF (480°C). Dimensional changes during this part of thetreatment are very small, so that in many cases the parts can be completely finished beforehardening.

Maraging steels can be welded without preheat in both the annealed and fully heat-treatedconditions. These alloys possess resistances to hydrogen embrittlemenl and stress-corrosioncracking that are generally superior to those of high-strength, low-alloy steels. However,the primary attribute of maraging steels and the main reason for the interest in them asconstructional alloys, particularly in extreme-duty applications requiring high strength-to-weight ratios, is their excellent toughness at high strength levels. In particular, theyexhibit markedly higher resistance to iow-stress fracture, even in relatively thick sections,than do conventional 0.3 to 0.5 %C quenched-and-tempered low-alloy steels, as typifiedby-AlSl 4340. The plane-strain fracture toughness (Klc) of the 18",,Ni maraging Mcels, atyield strength levels of 240,000 to 280,000 psi (1650 to 1850 MN/'ra-), is more than twicethat of the best 4340-type steels. The ultimate tensile strength and fracture toughnessof various maraging steel grades are shown in Figure 5.3.

( I l l \ l I M I N I U N I M i H i t i l l N I R t \ t , l l l

1 ollowins: the development of niaragini; steels, considerable effort has been devoted toMiuhtne the struc'ural characteristics of alloys of this kind. The results have heet1

Utscus-eJ m previous reviews, the most recent and comprehensive of which are those inReference- - v and -"" •' Although the various phase transformations and age-hardeningreactions that occur ,n maraging steels will he examined here in the light of the latest^formation available, numerous excerpts from the review paper m I lof.-cn \5.x\ havebeen included in the following pages.

- ! ! K , : : t '•! - I / . • ' ( ' I / ( I f l - i i ' - ' h W S

I' ha- heen -hown ihat alloying of ferrite with nickel not only ensures u niartensitic-iructure. which increase- th.e strength of ihe matrix, but al.-«o reduce-- the solubility of"ian\ element- i ]'i, \ lo. A I. etc I in y-iron I • urlhormore. the presence ol nickel lowers:!IL resi-tance •" liie crystal kitnee to the movement of dislocations and reduces the energyof the interaction of dislocations with interstitial atom- [.vV]. It promotes stress relaxationand. as a re>ult. reduces the susceptibility of the steel to brittle failure.

The marten-mc structure creates favourable conditions for uniform nucleation anddistribution of mtermetailic phases during ageing, thus ensuring higher plasticity andJiictilitv. Detailed -Indies of Fe-Ni martensites containing addition dements such asVi. Be. A I. Mo. Mn. Nb. Zr. \V and Cu. have shown that age-hardening occurs in thet'dii in 12(10 |- i?50 ni (,50 Ci range. However, the degree of hardening may vary appre-ciably with changes in the nickel content [5./f>]. or with variations in the ageing or reversionreacnon kinetic- Moreover, rather strong interactions can occur between specificcombinations of elements, such as cobalt and molybdenum. Nevertheless, the alloying.•lemenis can be classiiied qualitatively as " s t rong" (Be. Til. "modera te" lAI. Nb.Mn. Mo. Si. la. V and Wi. or " weak " (Co. Cu. 7s) hardeners.

•\s will be shown later, lhe main elements involved in the hardening of maraging steelsthrough formation of intermetaihc compounds, are titanium, molybdenum and. indirectly,cobalt. In mo->l ca>es nuiniitm plavs the double role of hardener and refining agent to lieup residual carbon. During solidification of the steels, this element tends to segregate orto precipitate in the austenite grain boundaries in the form of a network of Ti(C,N> carbo-nitrides. causing anisoiropy of the plasticity and reducing ductility [5.11]. Precipitationof compounds of titanium with carbon, nitrogen and oxygen was also detected on ageingmaraging steels between 700 and 900 F (370 and 480 C) [5.12\. The critical titaniumcontent above which ductility is lowered h.ith before and after ageing has been reportedas 0.8 to 1.2",, [5.5*]. It is also known that titanium additions lower the fracture toughnessof maraging steels (Fig. 5.3).

Molybdenum lowers the diffusion coefficients of a number of elements in the grain boundariesand thus reduces preferential precipitation of second-phase particles during ageing, therebyraising the ductility and plasticity of the aged steels. However, molybdenum also tendsto segregate during solidification, which again induces anisotropy of the plasticity andductility. Aluminium leads to limited hardening of martensile: at concentrations over0.2 to 0.3",, it lowers ductility both before and after ageing. Manganese additions leadto the formation of a martensitie structure at relatively low nickel contents, but have adetrimental effect on ductility after ageing. Silicon at concentrations above 0.1 ",;, reducesthe plasticity of the steels.

Some elements, while not inducing ageing of Fe-Ni martensite, increase the amount ofhardening that occurs on ageing by lowering the solubility of the hardeners in a.-iron.This is the case for cobalt, chromium and silicon, which respectively lower the solubilityof molybdenum and tungsten, of titanium, and of molybdenum and titanium. The decreasein the solubility of the hardener elements leads to an increase in the volume percentage

52

5. Ni-Co-Mo MANAMNG STfifiLS — PHYSICAL MFTAU.UKfiY

of ihc precipitates formed and reduces ihe work of formation of nuclei of the precipitatinephase, which, in mm. leads 10 an increase in she number of nuclei capable of growth at agiven temperature and. consequently, lo a reduction in the average distance betweenthese particles: each of these factors affects hardening.

In the same w;;y as nickel, cobalt lowers both the resistance of the lattice lo dislocationmovements and the energy of interaction between dislocations and interstitial atoms [5.9],furthermore, as cobalt raises the A/» temperature, increased amounts of alloying element*that ir.duce hardening during ageing can he added without leadinu to the formation ofresidual ULislcnitc following hardening. As regards chromium, this element increases thecorrosion resistance of the steels and raises the strain-hardening coefficient of murtensite.

In concluding this section, mention will be made of an empirical relationship | .\/ .f] lopredict the 0.2",,-otlsel yield strength in terms of ihe composition of maraging steels :

0.2",, Y.S. (in 10-' psil0 2",, Y.S. (in MN m=)

15.1104

(",,Co) - 28..1 <",,.V1o) - 80.1 (",,Ti|552(",,Tii.

This relationship applies only to IS",,Ni alloys and is valid up to about 300.000 psi(2100 MN m : ) : the values derived from it are accurate lo : 35.000 psi ( : 240 MN m ; i .independently of the method of preparation. The equation tends to confirm that thestrength of lXn,|Ni maraging steels results mainly from the effects of molybdenum andcobalt and. to a lesser extent, of titanium: molybdenum contributes about 50",, andcobalt 30",, to the tensile strength of these steels.

5.1.2. Conipttsiiitms

The compositions of current Ni-Co-Mo maraging steels are shown in Table 5.1. Asregards the lS",,Ni type, three grades were initially developed: these were arbitrarilyidentified by their typical yield strength '.allies expressed in I(V psi. viz. (200). (250) and(100). The strength levels attained depend primarily, on the combination of cobalt andmolybdenum, although increasing amounts of titanium are used in these three compositionsas a supplemental hardener. The excellent properties of these wrought grades promptedihe development of a cast version of this type of steel. As a result of a statistical studyof compositional variations aimed at obtaining a good compromise between tensile

TABLE S.I. — NOMINAL COMPOSITIONS OF Ni-Co-Mo MARAGING STEELS*

Alloy

18NK200)

18NU25O)

18NU3OO)

lSNi(casll

15Ni-9Co

l2Ni-2Mn

l8Ni(35O)

13NK400)

8Ni(500)

. l5Ni-15Co

IN-763 **

Ni

17-19

17-19

18-19

17

15

12

17.5-18.5

13

8

15

18

Co

8-9

7-8.5

8.5-9.5

10

9

12-12.5

15-lh

18

15

15

Mo

3-3.5

4.6-5.2

4.6-5.2

4.6

5.0

4.0

3.8-4.6

10

14

1.0

3.0

Al

0.05-0.15

0.05-0.15

0.05-0.15

0.1

0.7

0.1

0.10-0.15

—

—

0 0 5

Ti

0.15-0.25

0.3 -0.5

0.5 -0.8

0.3

0.7

0.2

1.4 - 1.7

0.2

0.2

0.4

0.05

OtherYear

announced

! l9hO__

_

—

—

2Mn

—

—

—

0.5V

1960

1960

1963

1963

I96h

1968

1968

196K

1971

1971

In wl.'\i, bal. Fc. The steels also containStilt in ihe product developmental stage.

- 0.03C. -• O.ISi, - O.IMn. 0.0IS. • 0.01P.

53

t i | i \ l I c H M \ I M \ i . H H . 11 M K I M . I I I M i l l s

strength and toughness, the lN\i(castl grade was developed [?.I4]. It is interesting tonote that iIK cobalt content of the >teel h;is been niised to 10",,.

The !5\i-l>Co steel I?.!?] and. more recently, the l5Ni-l5Co steel [5.Id] were developed!o pi oxide greater resistance to austenite reversion, which properly is of particular interestfor magnetic applications at temperatures up 10 800 I (425 C). As regards the steelJesiaiuued l2Ni-2Mn. it IN characterized b\ a 2",,Mu addition which permits the obtentionof a satisfactory nuirtensiiic >trucmre with a relatively low nickel content [5./").

Research aimed at developing maragmg steels uith higher strength led later u> arebalanced ls",,Ni composition having a nominal 350.(100 psi (2400 MN m-1 yield strengthand acceptable toughness [.\/iV]. This was achieved main!) by raising the Ti content.The increased cobalt content also contributes to strength and is effective in maintainingthe benefits of the unique cobalt-molybdenum interaction at lower molybdenum contents.The cobalt content in this high-titanium alloy is kept preferably to less than I5"o, so asto retain sufficient notched tensile strength. Reversion tendencies are reduced as a resultof the lower molybdenum and nickel contents. On the other hand, the l3Ni(4OO) [5.19]and NNi(500i grades resulted From an exploratory programme [5-2D] designed to determinewhether higher strength levels could be produced by heat treatment alone in Fe-Ni maragingsystems with modified compositions. As shown in Table 5.1. nickel, cobalt and molyb-denum are still the principal alloying elements in these steels: however, the nickel contenthas been reduced whereas the cobalt and molybdenum contents have been considerablyincreased. Titanium is still used in miner amounts but aluminium is not added. The13Ni(400) steel is of particular interest in that it offers a strength advantage over existingmaraging steels while maintaining a reasonable degree of toughness (<;/. Fig. 5.3). Finally.a composition aimed at reducing niicrosegregation effects and designated LN-763 is underdevelopment; it is characterized by a higher cobalt content, relatively low molybdenumand titanium contents, and a small vanadium addition [5.21].

5.2. Martensitic Transformation

5.2.1- Formation and Morphology of Mar tensile

Several of the basic characteristics of maraging steels are directly related *o the featuresof the Fe-rich end of the Fe-Ni phase diagram. According to the equilibrium diagram,the low-temperature equilibrium phases in iron-rich alloys are ferrite and austenite.However, on cooling an alloy containing around 10 to 25%Ni from the austenitic field,the austenite will not decompose into the equilibrium austenite and ferrite compositions,even if held for very long times in the two-phase region. Instead, with further cooling, theaustenite transforms to b.c.c. martensite by diffusionless shear, as in conventional steels(see Section 2.3.1). A general thermodynamic treatment of tht martensitic transformationin Fe - 9.5 to 33.2 at. "„ Ni has been developed [5.22]. The martensitic transformationtemperatures are shown as a function of nickel content in the metastable equilibriumdiagram of Figure 5.1. As stated earlier, the transformation exhibits a thermal hysteresiswhich increases with increasing nicke! content. Also, the Ms temperature decreases withincreasing nickel content: as an example, raising the nickel content from 20 to 25%decreases the M, temperature from roughly 390 to 100uF (200 to 40:C).

When the martensite is reheated, one of two things may happen. On the one hand, if the.alloy is brought to a temperature below the A, point, the martensite will decompose intothe equilibrium austenite and ferrite compositions, i.e., the martensite reverts to theequilibrium structures. The rate of this reversion reaction depends upon the temperatureand, fortunately for maraging steels, the rate at temperatures of the order of 900°F(480'C) is slow enough for considerable precipitation hardening to be achieved berorethe reversion reaction predominates. If, on the other hand, the alloy is heated above As,

54

Ni-f.'ii-Mo MANAGING STFFU.S — PHYSICAL METALLURGY

lrig. :i.4. — Structure nf l-'e-Ni-Mo alKi>«i.After J. Hoimcim et al. |.v-'.v],• q i iunchci l f r om 21110' |- 11 KlO't ' j* quenched f rom 21401- (12001:')• quenched f rom 22SO l r (1250 C'l.

301

S20

CD

FERRITE

M+F /

/ MARTENSITEx\

\ \oL0

\

5 iQ 15 20NICKEL CONTENT, w t . %

25

[lie mariensite transforms by a shear reaction back to an austenite of the same composition.In practice, even with relatively fast heating rates, some reversion occurs during heatingwhich influences the subsequent shear reaction.

Figure 5.1 indicates that only alloys containing up to about 33",,Ni will transform mar-tensitically. This diagram is actually rather oversimplified, because recent work hasrevealed that a surprising variety of transformations can take place in these alloys [5.23to 5.25]. In fact, three distinct structures, and hence three separate morphologies, havebeen shown to exist for the b.c.c. y. phiise obtained on cooling the f.c.c. y phase. In 0 to5",,Ni alloys, independently of cooling rale, and in 5 to 10"(JNi alloys for slow coolingrates, a structure comprised of equiaxed x-ferrite grains is obtained; these grains are lessregular in shape than is the case for annealed metal. A; sufficiently high cooling ralesmartensite will form in the 5 to !0%Ni alloys [5.26]. Increasing nickel contents lower thecooling rate necessary to form martensite, and at about 10%Ni a completely martensiticstructure is formed even with very slow cooiing: a typical example of this •* lath "martensite is shown in Figure 2.10 (p. 10). Lath martensite is found over the compositionrange from about 10 to 25%Ni. Finally, alloys containing more than 25%Ni transformbelow room temperature to a twinned martensitic structure. All these limits are notexact and vary, not only with cooling rate, but also with annealing temperature [5.27]and interstitial element content [5.23].

Figure 5.4 shows the composition range in which lath martensite forms in the Fe-Ni-Mosystem on quenching. In the case of maraging-type compositions, examination of a seriesof Fe-7Co-5Mo-0.4Ti alloys containing various amounts of nickel revealed that lathmartensite apparently forms for nickel contents up to 23%, whereas twinned martensiteis formed at higher nickel contents [5.29]. The lack of visible surface shears on prepolishedsamples has raised doubts as to whether the 18Ni(250) steel always transforms to lathmartensite [5..W], However, all the transmission microscopy evidence indicates that, withthe exception of the 25%Ni steel, all maraging steels normally have lath martensitematrix structures.

It is also known that the lath martensitic structure can be obtained in molybdenum-containing maraging steels by partial substitution of manganese for nickel. In fact,manganese may replace nickel in the proportion of 1 to 3, to a maximum allowable contentof 6%, without decreasing the Ms point below room temperature [5.17]. However, becauseof its embrittling effect, manganese should be kept to much lower levels to retain adequatetoughness.

55

I lUIALI l O M M M N i . H l l . I I S l R l N l . n l STEELS

5.2.2. Factors Controlling Unit Mar I ensile Formation

Ahhouch further studies are necessary to ascertain the precise conditions that determinewhether lath or twinned martensitc is formed, the two factors which seem to heimportant are the Af, transformation temperature and the stacking-fault energy (SFE)of the alloy. Lowering the ,W< temperature or raising the SFE should favour the formationof twinned martenshe. as was pointed out in Chapter 2 (Section 2.3.5). The SFE canonly be varied by adding an appropriate alloying element, whereas the Mf temperatuream be modified by both alloying and prior deformation. In some cases, the effects due to\t,< temperature and SFE may counteract each other.

The alloying elements commonly used generally lower the Ms temperature, but often theeffect of any individual addition is not constant and depends upon the total compositionof the alien. The effects of the solute elements commonly found in maraging steels on theUs and As temperatures of typical base compositions are shown in Figure 5.5. It can beseen that nickel generally lowers A/.,: in the case of Fe-Co-Mo compositions, this effectis greatest for nickel contents above 17.5",, [5.18]. Molybdenum strongly depresses theMs temperature, especially in steels with high nickel and cobalt contents. Chromiumalso causes a pronounced drop in the Ms temperature of Fe-Ni ailoys (cf. Chapter 8.Table 8.3): since the addition of chromium to a steel lowers both its SFE and its A/stemperature, this element exerts opposing effects as regards lath martensite formation.Increasing amounts of titanium, niobium, vanadium and silicon in a Fe-22.5%Nicomposition cause the A/,< temperature first to increase and then to decrease. Theprobable lowering effect of titanium on the Mx temperature of a Fe-Ni-Co-Mo maragingcomposition is also presented in Figure 5.5; it is likely that some interaction occurs betweenmolybdenum and titanium when co-present, as has been reported by one investigator[5J2]. Finally, increasing the aluminium content of the Fe-22.5",,Ni alloy gives rise toa slight initial increase in the Ms temperature with little subsequent effect [5.31].

700,

18 22 " 26Ni CONTENT, wt.%

1 2 3 i 5 6 0 0.5 1 15 t iMo 03NTENT,wt.% Ti.NbCONTENT,wt.%

1 " 2 3V,AI,Si CONTENT,wt.%

Fig. 5.5. — Effect of solute element contents on M, and A, temperatures of Fe-base alloys.• After G.W. TUFFNELL and R.L. CAIRNS [5.IS]. a After CM. HAMMOND [5.32],o After R.B.G. YEO [5.31]. ,-, After J. MANFNC ei al. [5.33].

56

5. Ni-Co-Mo MARAGING STEELS — PHYSICAL. METALLURGY

One rather significant exception to the above-mentioned trends is cobalt which, atmoderate levels, raises both the Ms and As temperatures of Fe-Ni alloys (cf. Chapter 2,Fig. 2.9, p. 9) and maraging steels in general \5.3I}. In practical terms, this has provedvery helpful in allowing higher alloy contents to be included in the steels while stillensuring that lath martensite transformation remains possible. As regards the effectof prior deformation on the Ms point and the resulting martensite structure and strength,this has been investigated in the case of low-alloy steels [5.34]. The observed loweringeffect on A/., probably also holds for maraging steels.

As discussed in Chapter 2 (Section 2.3.6). the general dependence of lath martensiteformation on SFE and Ms temperature factors is of interest as regards maraging steels,because there is some evidence that, after ageing, a lath martensiie matrix gives bettertoughness than does twinned martensite. Moreover, control of the M« temperature ofmaraging steels is also of importance to other properties. For example, their cold-workability is considerably improved on raising the nickel content from 18 to 24% andadding 3%Cr [5.35]. This depresses the Ms temperature of the steel below room tem-perature, while maintaining its M,i temperature above room temperature. This type ofsteel therefore possesses a metastable austenitic structure which transforms to martensiteon cold working, thereby increasing the strain-hardening rate.

5.3. Ageing of Martensite

The hardening induced in maraging steels during ageing may result from the followingtwo mechanisms :

— the fine, uniform precipitation of various intermetallic compounds or of austenite;as in the decomposition of other solid solutions, precipitation of stable compounds inFe-Ni martensites may be preceded by the formation of intermediate metastable phases:

— an ordering reaction in the cobalt-containing solid solution.

The extent of strengthening during ageing depends on the relation between the hardeningand softening factors. Foremost among the former is the distortion of the martensitecrystal lattice resulting mainly from an increase in the volume fraction of the precipitates.The softening factors include coalescence of precipitates, possible changes in the martensitesubstructure, a.id isothermal formation of austenite, which has been shown to occur attemperatures below the reverse martensitic transformation point. Many investigatorshave noted that some dislocation rearrangement, and usually a decrease in dislocationdensity, take place during the early stages of ageing; these changes are presumably partof the recovery reaction within the martensitic matrix. The onset of precipitation probablyprevents further dislocation motion.

5.3.1. Precipitation Reactions

In recent years, a considerable amount of effort has been directed towards both identifyingthe phases precipitated during ageing of martensite, and determining the shape, sizeand distribution of these precipitates. For this purpose, three-component and morecomplex steels have been investigated, mainly by electron microscopy and X-ray diffractionanalysis, although several additional identification techniques have been used. Forexample, in some cases the chemical composition of the extracted phases has beendetermined by microprobe analysis. Also notable in recent studies are the use of Mossbauerspeetroscopy and neutron diffraction analysis, as well as the analysis of the angles betweendiffraction spots. The influence of ageing temperature or alloying on various physicalcharacteristics such as lattice constant, electrical resistivity, dilatation and Young's modulushas also been studied in some detail.

57

C OKU 1 ( O N I U N I N U HK.HSIKE.Nf.iH STUKI.S

I ABLfc 5.:. CRYSTALLOGK.i'HlC DATA OF PHASES IDENTIFIED IN MARAGING STEELSAFTER W.B. HIASSON [!.-!?]

Phase

NiiMor-NuTi!V;Mi-

Ke-Ti

T-FeMo

-r-reFi

Structure ivpe

Orthorhombic Cu*Ti-type

Dt>;4 ordered hexagonal

Hexagonal C!4-i>pc

Hexagonal CI4-iypc

Tetragonal

Cubic CsO-ixpe

Hexagonal

Lattice

" it

« 5.0M.

2.5505 1

4.74 !

4.S13

sf.2118 !

4.146 j

parameter-,

c

b 4.224, <

!O<)f>7

7.73

7.855

4.SI3

25.78

IA)

<• a

4.448

3.2St)9

O.ttlJ

1.633

0.522

5.432

The results of the phase-identification studies have been ably summarized in References5.S and 5.9. A study of the tables presented in these papers reveals a number ofdiscrepancies, even as regards the precipitaies identified in alloys of almost identicalcompositions. This can no doubt be explained in part by the uncertainties inherent in thetechniques employed. The phases that have been identified are : NijMo [5.36 to 5.41],r.-NijTi [5.29.'5.41 to 5.43]. Fe:Mo [5.29, 5.37, 5.42. 5.44], Fe,Ti [5.38], a-FeMo [5.40,5.42. 5.44 to 5.46]. rr-FeTi [5.36], and u-Fe7Mon $.19, 5.44. 5.46]. Data on theirerystallographic structures are presented in Tabie 5.2.; The shape of the precipitates hasbeen variously reported .o be spherical [5.36. 5.3S, 5.43. 5.48], disk-like [5.48], ribbon-like[5.36. 5.3V] or needle-like [5.29. 5.36. 5.43].

Of the molybdenum-containing inte/metallic compounds, the one reported most often asprecipitating in 18",,Ni steels is Ni^Mo. The precipitated particles, which are rod-shaped,are abou'. 25 A wide and 500 A long in the peak hardness condition, with their longeraxes parallel to the 111 directions of the matrix [5.41]. The orientation relationshipbetween the NijMo orthorhombic precipitate and the b.c.c. martensite matrix has beententatively described [5.41] as :

(0»0)Ni,Mo // (0" )* ll00]N i j M o /

' 50.000Fig. 5.6. — Electron micrograph of I8Ni(250) steel aged 8 hours at 900°F (4803C).After J.M. CHILTON and C.J. BARTON [5.36].

5.357= compression

Lb%//COHla3aa%

compression

DOO]p//[iTi]a

186% expansion

a) atomic mislil alongcrystal axes of precipitate !p).

Fig. 5.7.— Distorted matrix lattice surrounding N13M0 precipitate in 18Ni(3OO) raaraging steel.After K. SHIMIZU and H. OKAMOTO 15.41].

b) distortion in (Oil) plane of matrix (1).

58

5. N i - C o - M o MAKACiING STF.f.LS PHYSICAL METAI I .URtiY

Fig. 5.8. - Electron micrograph of I3NK400) sleel aged•1 hours at 900°F <480"'C). After A. MAONEE ei at. [5.44].

35.000

i.e., the closest-packed plane and direction in the precipitates are parallel to those in thematrix. Figure 5.6 shows the distribution and morphology of the precipitate in thelRNi(25O) steel aged for 8 hours at 900'F (480'C), while Figure 5.7 illustrates schematicallythe distorted matrix surrounding an Ni3Mo particle.

It has also been suggested that the Ni.-|Mo precipitates produced after a conventionalmaraging heat treatment of several hours at 900°F are metastable, and that after ageingfor longer times and/or at. higher temperatures these precipitates are replaced either byFe2Mo [5.37] or a or phase [5.49]. The better lattice fit between Ni3Mo and the b.c.c.rnartensite matrix would initially favour precipitation of this metastable compound, butits growth would be limited by increasing coherency stresses which would favour thenucieation of the equilibrium precipitate [5.50].

Another tendency is to identify the molybdenum-containing precipitate as an Fe^MoLaves, [i-Fe7Moe or c-FeMo phase. Precipitation of such compounds, or at least ofthose of the AiB or A7Be types, is apparently favoured by lower nickel or nickel -+- cobaltcontents, or higher molybdenum contents [5.IG]; this effect might be related to a lowerelectron : atom ratio. The c-FeMo phase has been identified in the 13Ni(400) gradeaged af 900°F for 4 hours (Fig. 5.8).

As regards titanium-containing precipitates, they have been identified essentially in Ti-richcompositions; Y]-NijTi is the compound quoted most often, although a cr-phase hasoccasionally been reported. It is also possible that some of the titanium combines withmolybdenum to form a precipitate such as Ni3(Mo/Ti). The exact role of titanium duringmaraging of the Mo-containing ailoys is therefore rather difficult to assess.

Chemical analysis of precipitates extracted from I8%Ni alloys has shown that the truecompositions of the precipitates are not as simple as those represented by the formulaein Table 5.2 [5.38, 5.49, 5.51].. Th°v generally contain a considerable amount of iron and,in some cases, small amounts of other alloying elements. In particular, the nickel-molybdenum precipitate has been shown to contain too little nickel and too muchmoVodenum for it to be NijMo, and it has been tentatively identified by some authorsas Ni2FeMo [5.40]. Cobalt is usually present only in small quantities, if at all [5.51].A Mossbauer study [5,37] confirmed that no significant precipitation of cobalt occurredduring maraging.

Some evidence has been found that suggests that pre-precipitate zones may form duringthe initial stages of age hardening. From observations of diffraction streaks in a sampleof the titanium-rich 25%Ni steel after cold working, refrigeration and ageing for2 minutes at 890°F (475°C), it was suggested that there is a G.P.-zone stage in whichsegregates form parallel to the < l l l> a directions [5.43]. Diffraction pattern streakingwas also observed in the I8Ni(250) steel and some higher-titanium alloys [5.52], and thefollowing hardening mechanism involving titanium was proposed : lath martensite —>•metastable b.c.c. ordered Ni3Ti zones (D03 structure) - v Widmanstatten precipitation.

59,,

n M M M M . IIK.II SIP.lNdl H s n - H S

of M;I We T-NhTi pha>e (DO-M structure). In the same irivcs»!i$::!!'on, observation of>ever.il additional diffraction lines let! the authors to put forward another sequence.insolvini; molybdenum : lath martensite —* disk-shaped f.c.c. /»>ne> Uit, — 4.1 Al -••snheroidal le-Mo preeirtitaie. As regard.-, precipitation of the i -h tMo compound, workon Ke - !3Cr - 10 to IiiCo - 2 to 5\lo steeK [5.53] indicated the following sequence :formation of chromium-rich /ones —• intermediate precipitate — >• stable T precipitate.

Precipitates formed on ageing quite often appear to have been nucleated al dislocationsor at 'he ".v.'.r'.onsu-j !:iih boundaries. In general, ihe precipitates are uniformly distributed,and precipiu''---free /ones or coarse precipitates at grain boundaries are noi normallyloitnd. No d»mht the kith martcn.siie dislocalion structure is helpful in providing amia-i-uniform distribution of sites for precipitation. As aheady indicated, the precipitatesarc freo,uentl\ frehe\ed to be coherent with the matrix, and in many cases strain fields havebeen detected around the panicles \5.3n. 5.43. 5.4f>. 5.4V. 5.52]. An interesting observation!•» that a preferred precipitate orientation could arise because of the preferred orientationof the dislocations within the martensite matrix [5.54]. If precipitation occurs along theleneth of the dislocations, the precipitates would tend to lie in the 111 i directionsalready referred to [5.43].

5.3.2. Ordering

As already pointed out. an ordering reaction in the cobalt-containing solid solution cancontribute to the hardening of maraging steels on ageing [5.6. 5.4J. 5.32, 5.55. 5.56]. Asregards the binary systems involved in this type of steel, ordering has been observed overa wide range of solid-solution compositions in h.c.c. Fe-Co alloys; it is known to occurin fee. Fe-Ni solutions at 50 and 75 at.",, Ni. whereas no ordering has been observed infee. Co-Ni sohd solutions.

On the basis of neutron diffraction experiments, it was stated that B2-lype long-rangeordering developed on ageing an Fe-22.7Ni-19.3Co alloy [5.55]. However, neutrondiffraction experiments have also shown the absence of long-range order in ISNi - 8 W12Co maraging steels [5.49. 5.57]. From similar work performed on the l8Ni(35O) steJaged 3 hours at 950 F (510 C) [5.57]. it was suggested that a high degree of short-rangeorder is established in localized regions. Since the magnitude of the interatomic attractionsdecreases in the order Fe-Co. Fe-Ni and Co-Ni, Co-Fe-rich short-range ordered regionswould form preferentially, the remaining regions being nickel-rich.

These short-range solid-solution atomic arrangements may control the processes leadingto precipitation of the Mo-Ni and Ti-Ni intermetallics in the nickel-rich regions. Thediffusion of iron towards these regions being inhibited by the Co-Fe attractions, longerlimes or higher temperatures are required for the replacement of the nickel-rich phasesinitially formed by iron-rich precipitates [5.58]. This reasoning is compatible with themetastability of the phases precipitated in maraging steels on ageing at 900°F (480°C)[5.37].

5.3.3. The Cobalt!Molybdenum Interaction

Age hardening can be produced in ternary Fe-Ni-Co alloys, and with high nickel andcobalt contents very high strengths can be achieved. Figure 5.9 gives the curves of yieldstrength versus cobalt content for two different Ni levels and two different maragingheat treatments [5.21]. The strength appears to increase approximately linearly withincreasing Co concentration. However, one of the most important contributions to thestrengthening of maraging steels is that due to the combination of cobalt and molybdenum.As was evident from Figure 5.2, these elements are potent hardening agents .when addedsimultaneously to iron-nickel alloys. In fact, the hardening obtained when Co and Mo

60

5. Ni-Co-Mo MARAGING STEELS — PHYSICAL METALLURGY

are present together was found to be much greater than the sum of 'he strength :ncrementsproduced when these elements are added individually [5.49, 5.59 to 5.61],

A comparison of the strengths of a number of ternary Fe-l8Ni-X and quaternary Fe-l8Ni-8Co-X alloys (where X --- Al, Be, Mil, Mo, Nb, Si or Ti1 [5.59] has shown that, with theexception of the molybdenum-containing alloy, the addition of 8"/,,Co increases theyield strength by about 20.000 to 45,000 psi (140 to 310 MN/m2). This is of the order ofthe strength increase produced by cobalt alone, as illustrated in Figure 5.9. Thus, in thesealloys, cobalt supplies only a relatively small, roughly additive strengthening contribution.In the molybdenum-containing quaternary alloys, however, the strength was increasedby up to 75.000 psi (520 MN/m-) on adding cobalt.

The part played by cobalt is not clear, and several mechanisms have been proposed toexplain the enhancing effect of this element on the hardening of Fe-Ni-Mo and Fe-Ni-Cr-Mo matrices caused by molybdenum.

Examination by transmission electron microscopy has led to the suggestion that additionof cobalt results in a finer dispersion of precipitates in the molybdenum-bearing alloys.Since cobalt has not been found to be present to any extent in the molybdenum precipitates[5.51], it is generally considered that cobalt may lower the solubility of molybdenum inthe martensite matrix and thus favour precipitation of molybdenum compounds [5.46,5.52, 5.59]. Electrical-resistivity measurements have also shown that the principal effectof cobalt is to increase the amount of molybdenum precipitated. This technique provedvery useful since it was found that resistivity is directly proportional to the amount ofmolybdenum in solid solution, while changes in cobalt content have a negligible effecton resistivity [5.60]. The measurements also revealed that the resistivity changes due tothe recovery of the matrix during ageing were very minor compared with those producedby precipitation. Figure 5.10 illustrates the resistivity changes observed in Fe-18Ni-5Moternary and Fe-l8Ni-5Mo-8Co quaternary alloys during ageing at 850°F (455=C). Bothcurves show an initial loss in resistivity due to precipitation of molybdenum, followedby an upward turn due to austenite formation. The cobalt addition tends to acceleratethe ageing reaction and to remove more molybdenum from solution.

It has also been suggested that cobalt may alter the dislocation structure of the martensitematrix to provide more numerous and more uniformly distributed nucleation sites for

- 2 0 0 0

• t 0 - ••• : •• •'_ •-. - v - B

COBALT CONTENT, w t . 7 . -

Fig. 5.9.— Yield strength of Fe-Ni-Mo alloys. After S. FLOREEN [5.21].• Maraged 24 hours at 800°F (425°C)• Maraged 3 hours at 900°F (480°C).

40Anneale .001 .01 ; o.i l.o

. AGEING TIME, hours1000

Fig. 5.10. — Electrical resistivity of two Fe-18Nialloys as a function of ageing time at 850°F (455°C).

After D.T. PETERS and C.R. CUPP [5.60].

61

, HIlill-SMilMilH Stalls

subsequent precipitation [5.46. 5.51\. Moreover, considerable refinement of the atistenitearains occurs and the refilling inanensitic structure is therefore much finer [5.-J6]. Cobaltmay also decrease the SFL of the matrix in the ausienitic state, which v\otild discouragecross-slip and retard cell growth. The mean dislocation density of marlensite would heincreased, again providing more nuclcation sites for precipitation [5.45, 5.61].

While all these effects are possible, it must be remembered that, in quaternary Fe-lSNi-•SCo-X alloys, cobalt exerts a strong hardening effect only in conjunction with molybdenum.while with other hardeners cobalt product's on|\ a small additive strengthening effect[.\5'J]. For this reason, the change in the martensite dislocation structure due to cobaltwould seem to provide little hardening wr se [5.S].

5.3.4. Kinetics

As discussed in the previous sections, a number of reactions can take place duringmaraging. and the kinetics of these can be studied reasonably well before reversion of themartensite matrix to austenite becomes a major factor. Firstly, several authors haveobserved some dislocation rearrangement and loss during the initial stages of ageing.Internal-friciion experiments [5.62] have led to the conclusion that this recovery reactionis due to aihermal dislocation glide, and ends in precipitation.

As regards the precipitation-hardening process, it takes place very rapidly in most alloys[5.5rt]. Figure 5.11 illustrates the hardness changes observed in an Fe-18Ni-5Mo and anFe-l8Ni-SCo-5Mo alloy on ageing at 850 F (455 Q . It can be seen that the hardness ofboth alloys rises steeply after ageing for as little as 1 to 2 minutes. As slated in the previoussection, small but real changes in the resistivity of the same alloys are noted after even theshortest ageing times (Fig. 5.10). From such results, it has been concluded that theincubation lime for age-hardening is zero [5.6U. 5.62].

In many cases the isothermal ageing kinetics can be conveniently described by a relation-ship of the type : _J.Y v,, — Kt". where .v is either the hardness or the electrical resistivityat time /. .v,, is the value of this property in the annealed condition, and K and n areconstants [5.56. 5.60. 5.63]. Quite often the values of the time constant n are of the orderof 0.2-0.4. i.e.. distinctly less than the n value of 0.5 for the idealized case of diffusion-

.01 0.1 1.0AGEING TIME, hours

Ternary Quaternary

Coherent precipitation $ II

No coherent precipitation O D

_yl .V.. % , - , . . _ " :

iV #» . 15 20

at.% N i o r Ni+Co25 30

Fig. 5.11. — Hardness of :wo Fe-18Ni alloys asa function of ageing time at 850 F (455"C).After D.T. PETERS and C.R. CUPP [5.60].

Fig. 5.12. — Compositional limits for occurrence of cohe-rent precipitation in ternary Fe-Ni-Mo and quaternaryFe-Ni-Co-Mo rnartensites. After J. BOURGEOT et al. [5.28].

62

Ni-lo-M.) MARACHNC, STEELS PHYSICAL METALLURGY

controlled growth of platelets. Furlhermore. for all the I8°(,Ni-type alloys studied, thenominal values of the activation energies are fairly lov., being typically of the order of30-50 kcal/molc, which is well below those commonly observed for substitutjonal-elementdiffusion in ferrite.

On the basis of these results, it was generally concluded that the maraging kinetics areprimarily determined by the martensite matrix structure. The absence of an incubationperiod was attributed to the elimination of the free-energy barrier lo nuclcation underconditions of high supersaturation and precipitation at dislocations [5.60]. The lowvalues of the time constant n and the activation energy are commonly interpreted in termsof pipe diffusion through the dislocations present at a high level of density in the lathmanensite matrix. Evidence of the migration of substitutional solute atoms on dislocationsites in martensite emerges from (he observation of serrations on the s'ress-strain curvesfor Fe-Cr-Ni, Fe-Cr-Co and Fe-Co-Ni martensite alloys tested at the usual strain ratesof 10-3 t 0 io-6 s-i a n ( j at temperatures ii; the 66O-75OCF (350-400°C) range over whichageing may possibly occur [5.64].

However, other studies suggest that this interpretation is not completely correct. Theageing kinetics of an Fe-8Ni-13Mo alloy in the ferritic, cold-rolled ferritic, or martensiticcondition [5.65] were found to exhibit noticeable incubation times. Furthermore, theactivation eneny values for age-hardening of all three structures were of the order of65 kcal/mole. Similarly, a value of 60 kcal/mole was determined in the case of a martensiticstainless steel [5.56]. These results show that activation energies of the order of thosefor diffusion of substitutional alloying elements can be obtained during ageing in a lathmartensite matrix. In order to explain the rapid hardening of 18%Ni-type maragingalloys, it was accordingly suggested that the absence of incubation times is possibly relatedto the tendency of the higher-nickel alloys to form A3B precipitates and to the close fitwith the b.c.c. matrix that is generally possible with precipitates of this type [5.8]. Thetemperature dependence of nucleation may also be responsible for the low activationenergies observed.

As regards the effect of ageing temperature, it has been observed [5.60] that, contraryto the behaviour established for the Fe-18Ni-5Mo alloy, there is a pronounced discontinuityin the ageing kinetics of the quaternary Fe-18Ni-8Co-5Mo alloy at about 85OT (455=Q.This discontinuity shows up on Arrhenius plots of times to reach constant hardness orresistivity values, and in the peak hardness vs. ageing temperature curves. More recentstudies [5.28] on Fe-Ni-Mo and Fe-Ni-Co-Mo alloys have shown that two precipitationprocesses are operative on ageing the quaternary alloys. The first takes place within thematrix and is predominant when the ageing temperature is below 840GF (450°C), whilethe second occurs preferentially on dislocations, predominating when ageing is performedabove 840T. This complex precipitation behaviour had already been identified on thebasis of electrical resistivity measurements [5.67].

The hardening that results from holding at temperatures below 840°F is probably due tothe formation of ordered precipitates, 10 to 50 A in diameter, that are coherent with themartensitic matrix. They appear to have a hexagonal structure with a = 7.02 A andc = 2.48 A, and the following orientation relationships with the matrix :

[100]p // [211]a [011]P // [ l l l ] a

There would thus be four groups of orientation relationships with respect to the matrix.Figure 5.12 shows that the rather well-defined compositional limit for the occurrence ofcoherent precipitation appears to foilow approximately an electron concentration contourwhich corresponds to an average number of paired electrons per atom, epja, of about6.0. The formation of fine, coherent precipitates of an ordered phase on ageing the13Ni(400) maraging steel at temperatures between 750 and 840°F (400 and 450°C) has

63

U H I M I U I M M M M . MICH S I K K N C i l H S H U I S

tii.--.irun micrograph. 40.1HK)M HJectron microdifTryciion. show-

ing ordered hexagonal phase <•> Inicrprcuuan of diffraction pattern.

Fig. 5.I.V Microstrudure of I3NU4OO> alloy aged 1000 hours at 75OF 1400 C). Aftei A. MAGNI-K ei •;/. [5 7\.

also been reported [5.7]. The microslructure of she I3Ni(400) alloy aged for 1000 hoursat 750"F is shown in Figure 5.13 together with the corresponding diffraction pattern.the occurrence of such coherent precipitation is not entirely unexpected, since clusteringis known to occur in Fe-Mo supersaturated solutions [5.68 to 5:70].

The matrix precipitation in quaternary alloys has been attributed to the higher super-saturation of molybdenum in the presence of cobalt [5.60]. However, the similarity of theresults found for ternary Fe-Ni-Mo and quaternary Fe-Ni-Co-Mo alloys (Fig. 5.12)implies, on the contrary, that cobait does not play any special role in this particularprocess [5..iS']. On the other hand, calorimetric analysis [5.7/] and resistivity measurements[5.72] performed during ageing on the ferrous martensite of ternary and quaternary alloyscontaining chromium, nickel, cobalt and molybdenum suggest that the low-temperatureageing reactions in maraging steels are governed primarily by the formation of Ni- andCo-rich zones, and that the chromium and molybdenum only participate in an auxiliarymanner.

As regards the second mode of precipitation, i.e.. that occurring on dislocations on holdingat temperatures above 840 F. it ieads to relatively faster hardening of the martensite; thepresence of cobalt in the quaternary alloys accelerates the precipitation reaction in thesame way as does raising the molybdenum content, and confirms that the former elementincreases the supersaturation of the latter in these alloys as compared with ternaryalloys [5.28].

5.4. Austenite Reversion

As stated previously (</. Section 5.2.i), the Fe-Ni martensite matrix is metastable, andafter prolonged holding at elevated temperatures below ,4*, will eventually decompose by adiffusion-controlled reaction to ferrite and austenite. Depending on the temperatureand composition, the austenite formed may be so enriched in nickel that its Ms temperaturelies- well below room temperature and austenite remains in the microstructure onsubsequent cooling.

Potentiokinetic studies on ternary and quaternary Fe-l9Ni-base alloys with 5%Mo and9%Co additions and on the l8Ni(300) maraging steel subjected to different anisothermaltempers [5.73] have established that the formation of reverted austenite occurs througha diffusion mechanism followed by a rapid shearing process, or through the simultaneousaction of both mechanisms.

Ni-Co Mo MARAGING STF.F.l.S -- i'HYSir \1 . MFTAI.I.IJROt

The rate of austenile formation cs quite sensitive to composition. In binary Fc-N alloys.increasing nickel contents generally tend lo accelerate austenite formation " ?4],Additional alloying elements can also markedly affect the reversion reaction. The cii'cctof an individual alloying element on reversion may be partly explained on the basis ofwhether it tends to stabilize the anstcnite or the ferrile phase [5.75]. However, a muchstronger effect probably results from ihe change in matrix composition thai accompaniesprecipitation. Titanium, for example, has been found to retard reversion markedly, dueto the fart that formation of Ni3Ti lowers the nickel content of the matrix [5.74].Molybdenum additions, on the other hand, favour reversion [5.74. 5.76]. In this case.reversion is believed to be associated with the dissolution of Ni-,Mo and the formationof Fe2Mo [5.6.?, 5.74].

The effect of cobalt and molybdenum on reversion in maraging steels during temperinghas been established from a comparison of the dilatomelric behaviour of Fe-Ni, Fe-Ni-Co,.e-Ni-Mo and Fe-Ni-Co-Mo alloys [5.77]. It was found that, whereas the martensite-io-austenite transformation occurs in one step in the binary alloys, two successive stepsare required in the more complex alloys. As indicated in Figure 5.14, the initial rnartensitefirst decomposes into two distinct phases with different nickel contents. Next, the low-nickel ferritic phase transforms to either a single austenitic phase, aiso depleted in nickel(Fe-Ni-Mo and Fe-Ni-Co-Mo alloys), or to two different phases with different nickelcontents (Fe-Ni-Co alloys). Therefore, it is apparent that cobalt and molybdenumadditions each play a specific role, which is related to their effect on the \fs temperature.However, when both elements are added together the effect of molybdenum ispredominant.

Austenite usually starts to form at the martensite platelet boundaries; once the reactionhas progressed far enough, the structure is lamellar in appearance, with elongated austenileribbons strung out along the boundaries, as shown in Figure 5.15. In addition, smalleraustenite pools form within the platelets. The precipitated particles have occasionallybeen observed to dissolve in the austenite.

With prolonged holding at temperatures of the order of 750-1300'F (400-700X), veryappreciable amounts of stable austenite can be formed. For example, maximum amountsof stable austenite occur on holding the 18Ni(250) steel for 1 hour at 1110-1200°F(600-650°C) [5.78]. In general, the stable austenite produced by reversion cannot betransformed to martensite by refrigeration. Cold working, however, will cause it totransform [5.79]. The results of a recent investigation [5.80] suggest that austenite

SYMBOLS

a = ferriteY = austeniteM = martensiteC = nickel contentMs = martensite trans-

formation point

SUBSCRIPTS

p - Nl-depleted

r = Ni-enriched

o = initial

1st step MARTENSITE

Co

- * Ctp + Yr

Z \ on tooling. Yr—Mr

Fe-Ni-MoF*-NI-Co-Mo Fe-NI-Co

2nd_stet>7

C Y D< C °

on cooling. Yp-*M p

r.on cooling.Yr~*Mf

Fig. 5.14. — Schematic diagram showing possible paths of mar-tensite-to-austenite transformation in Fe - 20Ni - 3 to 15Cc,Fe - 20Ni - 4 to 8Mo and Fe - 20Ni - 5Mo - 2 to 9Co alloys.

After C. SERVANT and G. CIZERON [5.77].

Fig. 5.15. — Electron micrograph of13Ni(400) maraging steel agsd 4 hoursat 1110°F (600°C) and air cooled.After A. MAGNEE et at. [5.7]. x 30.000

i >U \1 , i i >\ I \ I M \ i . Mil .11 s i XI M . Ml S i l l 1 S

^ lii;. 5 U".-- LlTtvloftcmpi.'1'.iiBtcrnpcrii-^-'^ uiredinic; ~hlon saturation in.i^ncli/a-

i" non and C urie point ol 1SNI<-\^<M McelAf:«r 1' i.n.i sun; i* ~v).

L-. Ihe ausicniU', or '•; ph.ue. sLibilvcddown :o room (cmpctamre (in-crs inihc region marked A on the i';ft-hand siden! ihc figure. The IVW'.TI marked Hconespomis \o the v.nsiabih.'ed •• phase,uhich decoMinÂt'H to -i on coohnu,[he appro v..'3U!e tempcraUm.1'. to: tinsdc : ^r.'i'o.Mtion arc shown in the figure

stabilization is associated with a siopr--j;o in I he growth of the mariensite platelets bypinninfl of the austenite-martc-i.-site interface. During slow cooling, this blocking would bedue to the magnitude of ihe strain energy introduced by the early stages of the austenite-to-martcnsite transformation: on the other hand, during isothermal tempering, the inter-facia! energy vends to become preponderant because of the change of the mariensite,austenite interface from a partly coherent type to a totally non-coherent one.

The presence of austenite in the structure reuuecs the strength but considerably increasesthe uniform elongation during straining, which has been shown to he very helpful inoperations such as deep drawing [5.35. 5.79]. On the other hand, although the effect of<>\i.'rageing on stress-corrosion beh'.ivio r is not clear-cut, formation of large amounts ofreverted ausienite during an overagetng treatment seems to increase the stress-corrosionresistance of maraging steels [-5.-/-J.

The magnetic properties of maraging steels are also afieeted by the presence of revertedaustenite [5.16. 5.50. 5/>J. 5.74. 5.79]. As an illustration. Figure 5.16 shows the effect, oftempering temperature on the saturation magnetization and Curie temperature of thelis'Ni(250) steel. As austenite starts to form, the saturation magnetization \the value ofwhich depends inversely on the amount of austenite present) decreases appreciably.Simultaneously, there is a marked increase in coercive force. This is consistent with thefact that large changes in coercive force arc to he expected only when particles of thedimensions of the domain walls, i.e. ~ 1000 A in size, are present. Precipitate particlesare generally too small to affect ;he properties significantly, but the austenite paniclesformed by the reversion reaction are of the right size.

Austenite can also he formed by heating an alloy above its ,-J., temperature, but in thiscase the martensite transforms by a shear reaction to an austenite of the same composition.Studies of this reaction in Fe-Ni binary alloys have shown that the resultant austenitecontains a high density of dislocations [5.81 to 5.83]. Orientation relationships indicatethat the crystallographic nature of this transformation is similar to that of the austenite-to-martensite transformation. If the heating rate is low enough, some reversion of themartensite matrix to austenite • ferrite by the diffusion-controlled reaction may occur.When this happens, the subsequent shear reaction may be affected, apparently due to thefact that the initial reversion reaction tends to stabilize the martensite and to raise the As

temperature. This effect might be due to mechanical stabilization by the austenite particlesformed during heating, and/or to the corresponding decrease in the nickel content of themartensite. When additional alloying elements are present, the changes in the shearreaction with composition and heating rate can become very complex indeed. In thecase of the I8%Ni steels the general trend still appears to be that decreasing the heating

66

\ i ( o - M u \ I \ K \ ( , | \ ( , S i l l I S P H V S K - \ l . M L I A L L L ' k G Y

rate raises the .-(., temperature [5.W]. but the multiplicity of reactions makes inlerpretationof ihe data difficult.

As the temperature is increased well above .),<. a homogeni/ation reaction becomes pre-dominant. /.<•.. precipitales dissolve and concentration gradients due to precipitation orreversion are removed [5.7-/|. The lime and temperature required for completehomog.-nization depend, of course, upon the composition and history.

5.5. Strength-Toughness vs. Structure Relationship

The primary attribute of maraging steels is their excellent conihiiviiion of strength andtoughness. As discussed in Chapter 2 (Section 2.4.3). approximately one third to onehalf of the yield strength of fully heat-treated IS",,Ni maraging steels can be ascribedto the strength of the Fe-Ni lath martensite formed on cooling from the aiineaiiimtemperature, the rest being due to the line dispersion of precipitated particles which formon ageing [5.9]. It was also shown that about three-quarters of the overall strength ofFe-Ni lath martensitcs stems from solid-solution hardening, and the remainder is dueto the transformation substructure in the martensite (cf. Chapter 2. Fig. 2.13. p. 13).The reasons for the superior toughness of the I8"oNi maraging ste;ls are not clear. Ithas often been noted thai elimination of carbon and other deleterious impurities shouldbe beneficial in this respect, as should the relatively uniform precipitate distributionachieved in a lath martensite matrix by age-haraening. Molybdenum appears to playan important role in miniinmns einbrililement of the prior austenite grain boundaries:for instance, it has been reported that low-energy intergranular fractures of Fe-18Ni-XCo-base alloys hardened with aluminium or titanium can be eliminated by addition of2",,Mo [5.5V].

The cfleets of varying the nickel, cobalt, molybdenum and titanium contents on mechanicalproperties and microsegregation in maraging steels were recently examined [5.2/]. Anickel content of i8",, gave the best mechanical properties. Molybdenum additions raisedthe strength and gave substantially better impaci values at all strength levels. Increasingthe cobalt level from 12 to 20 "„ gave continuous increases in strength and decreases intoughness in Fe-l8Ni-Co base alloys. Good tensile and impact properties associated withreduced microsegregation were obtained on increasing the cobalt content from 8 to 15 °,,and lowering the molybdenum content from 5 to 3 % and the titanium content from0.4 to 0.1%.

The possible role of the high nickel content of the matrix in preventing fracture must alsobe considered. In low-alloy steels, nickel reduces the tendency to cleavage and lowers theductile/brittle transition temperature. In maraging steels, this effect of nickel could behelpful in minimizing hydrogen embrittlement, where quasi-cleavage fractures havegenerally been observed. Normally, however, fracture in maraging and other high-strengthsteels occurs by means of void nucleation giving rise to localized ductile fracture, and itis not certain that nickel would be helpful in preventing fractures of this type.

What is certain is that high nickel contents associated with low carbon and impuritycontents and a lath martensite matrix with a fairly uniform precipitate distributiondo not constitute sufficient factors to account for the toughness of the maraging steels :all these criteria are common to the whole family of maraging steels, yet the fracturetoughness of the 18Ni(25O) grade is significantly higher than that of the other gradescompared at equal strength levels [5.59]. Since the 18Ni(250^ grade is age-hardened by aMo-rich precipitate which is metastable and eventually £jes into solution, it has beensuggested [5.5] that this process may start during maraging to give a structure of Ni3Moparticles with adjoining austenite that might be more effective in preventing voidformation at the precipitates, and thus in delaying fracture.

67

I i I B M I 1 ( < M \ I M V . H k . l l S T R t N C i l H S T E E L S

,,. NWo-Mo MAR.\C;IM; s n i i s — THE COWEMIONAL GRADES

Although 'his chapter will deal essential!) with the three basic maragini: steel grades,|SNii2iH'ii. LsNii25Oi and lSNi(3OOi. data on some of (he other grades lisied in Chapter 5t Table 5.M will be included where appropriate. This will apply in particular to the castgrade (Sections (>. I. (v2.I, d.2.41. the heat-resisimg grade (Section 6.2.21 and the magneticgrade (Section 0.2.2).

(v i I'rimarx Processing

The wrought !*"..Ni steels can be prepared both h\ air and vacuum-induction melting.although fur the higher-titanium grades \acmim melting is preferable. Kemelting ina consumable-electrode furnace decreases segregation and improves cleanliness, leadingto superior properties. The cast grade, developed in 1%3 (see Chapter 5. Section 5.1.21.possesses sufficient castability and fluidity for its melting and casting to be performedm air ['>./]. Vacuum-induction melting of this steel has been recommended to prepareinvestment castings with good surface tinish [6.2. 6..*]. Finally, a recent study on theISNi(300i grade has shown that segregation-free products can be obtained from pre-alloyed powders b\ hot extrusion of canned billets [6.4].

The wrought maraging grades can be easily hot worked using standard procedures suchas rolling, forging (including drop forging!, drawing and extrusion. Prior homogenizationb> soaking the ingots at about 2300 F (1260 C) is recommended. Hot working maybe c.irried out between 2300 and 1500 F (1260 and 815 Ci. but scaling is minimized ifthe siarting temperature is limited to 2100 to 1900 F (1150 to 1040 Cl. Finishing at lowtemperature i- desirable for the obtention of uniform smaii grains and optimum mecha-nical properties. These steels have been shown to be self healing and to forge-weld likecarbon and low-alloy steels [rt.5. rt.ft].

Heal treatment of these steels generally comprises solution annealing at 1500 - 20 F(SI5 j_ 10C) for a minimum of 15 to 3.0 minutes for 0.05 in. (1.3 mm) thick sectionsand for 1 hour per inch for heavier sections, followed by air coofing [6.5]. At this stagethe material consists of a soft, low-carbon, high-nickel martensite which is readilyamenable to machining. Since the martensitic transformation is independent of thecooling rate, hardenability is not a problem. Maximum strengthening is obtained onageing at 900 •-• 20 F (480- 10 C) for 3 lo 6 hours, and air or furnace cooling.

As regards the cast grade, the optimum heat treatment consists in homogenizing at1800 F (980 C, for 4 hours, overageing at 1 !00'F (595=C) for 4 hours, solution annealingat IsOO F (815 Cl for 1 hour per inch, and ageing at 900F (480"C) for 3 hours [6.3].

6.2. Properties

6.2.1. Strength-Toughness Characteristics

A full background of information on the mechanical properties of the wrought maraginggrades mentioned in the preceding section will be found in [6.7]. Typical elastic, tensile,and toughness properties of these and the cast and sintered grades are summarized inTable 6.1. It is seen that the wrought and cast grades have good combinations ofstrength, ductility and toughness. As regards the powder-metallurgy l8Ni(3OO) steel,it is seen that its tensile properties are equivalent to those of the wrought product, whileits room-temperature Charpy V-notch toughness is higher; this improvement is attributedto the fact that use of segregation-free spherical powders produced by the rotatingelectrode method has eliminated banding [6.4].

68

6. Ni-Cu-Mo MAKAOINCi SIHfcLS THE CONVENTIONAL ORADES

TABLE 6.1. TYPICAL KOOM-TEMPERATURE MECHANICALPROPERTIES Ol THE CONVENTIONAL MARAGING GRADES

Material

ISNiCOO)bar '

ISNiCSO)bar '

ISNi(3OO)bar"

Casi grade'1

P M bar'

L!

IO1 psi

195-230

245-270

265-305

249

282

IS..

MWm-

1345-15X5

169(1. ,1M)

1X25-2105

17IX

1946

0.2'

10' psi

190-225

240-265

260-300

237

27S

,. VS..

1310-1550

1655-1X25

1790-2070

1635

I9IX

El.(2 in.).

6-12

6-10

5-10

10

10

R. A..

35-67

35-60

30-50

40

54

C\ 'Nimpact

fi.lh

26-50

18-33

12-19

i5

23

./

35-6X

24-45

16-26

20

31

K

|Q« j.-.in.

A/.Vni -! -

100-160

90-150

80 130

-

110 176

99-165

US-143

N.

L.

(A

(A

(A'

T.S.

T.S.

1.510)

1.5'Ol

1.5i 10)

-

Hard-ness.

R.

44-68

4K-5O

51-55

-

5"1

' "

Moduli.

<7 vIO"psl —

26.2 1X1( e l i t s t i c i n }

27.0 1X6(elasuein I10.4 72(rigidity)

27.5 190(elasticity!

— -

Pois-son'sratio

0.26

0.3

0.3

-

-

Ref.

[6.5]

\6.5]

[6.5]

[6.3]

V>-4]

Cuniiiiion : 1500UF (SIS' Ci - I h 900 c F (480^CI - 3 h.l imdilion : 1800 :K ("980 C.) - 4 h ' IIOff'F ( 5 9 5 ° O - 4 h - 1500°F IS15"O • 1 h ')00"F (480 CI - 3 h.Rolalinu-elcctrode processed !8Ni(3OOI powder extruded a! lfi5OcF (900"C! using a ram speed of 100 in. m : n . followed b> water quench ing : used at•Hid I 1480 Ci tor 3 h.

In Figure 6.1 the fracture toughness of the three basic maraging grades is comparedwilh that of M.eels of similar strength. The superiority of '.he former is immediatelyobvious.

It has been shown [6.9] that the puriiy of the raw materials used strongly affects thetoughness properties of the steels. In particular, the sulphur, carbon and nitrogen levelsshould be carefully controlled in order to prevent harmful grain-boundary precipitationof intermctallic c impounds [6.10]. This type of grain-boundary embriitlement can bereduced significantly tnroiijh small refining additions of magnesium [6.11]. Similarly,restricting the carbon level io less than 0 01"{ and the sulph.,: and n;( "ogen levels to lessthan 0.005"-,, greatly enhances transvc-: ductility \<,.li\. L. a more recent study [6.13]on the 18Ni(3OO) grade, ihe recommended carbon limit for optimum toughness has beenset at 0.005";,, while the simultaneous presence of siii-on and manganese at the 0.15°,,

Fig. 6.1. — Effect of strengthlevel on fracture toughnessofrepresentative high-strengthsteels. After J.C. HAMAKER

and A.M. BAYER [6.8].

1250ULTIMATE TENSILE STRENGTH, M N / m *

1500 1750 2000

ol_200 250 300

ULTIMATE TENSILE STRENGTH, 103 psi

69

t•ifflALT-l C M AININii HIliMSTRf-MiTH STFEl S

20 r

.300!—

280

-,2200

40 50V.COID WORK

40 " 50COLD WORK

Fiu. h.2. — Effect of cold work and maraging parameters on strengthfracture toughness of 18Ni(25O) crade. After E.P. GILEWICZ [6.

[nilial condition (0.375 in. plate) : 1500 F (815 Cl - 1 h, AC.

and18],

330

30 40 50 60V.COLD WORK

Fig. 6.3. — Effect of cold work and maraging parameters on strength andfracture toughness of 18Ni(3OO) grade. After E.P. GILEWICZ [6.18].

Initial condition (0.375 in. plate) : 15003F (8I5°C) - 1 h, A C.

70

(,. Ni-Co-Mo MARAG1NG STEELS — THE CONVENTIONAL GRADES

level has been reported to be highly detrimental to this property. 18/,,Ni maragingsteels do not appear to be sensitive to the elements phosphorus, antimony, arsenic, tin,lead, and bismuth, but exact limits arc not known;'chromium, copper and tungsten inamounts of 0.5",, appear to be harmless [6,14]. Vacuum-consumable electrode rcmeliinghas been found [6.15] to increase significantly' the short transverse strength of theI8NH2OO) grade, and to a lesser extent that of the l8Ni(250) and I8Ni(300) grades;moreover, the short transverse ductility of all three steels was considerably increased.Vacuum melting also improves fracture toughness appreciably [6.16], although someevidence to the contrary has been presented in the case of a 1.2%Ti composition [6.17],

As regards heat treatment, the early developmental work on the steels soon showedthat maximum ageing response was obtained on treating at 90CPF (480 C) (see. forinstance, Ref. 6.IS), and this temperature was thus adopted as the standard ageingtemperature. However, there are cases in which departure from the standard practicemay be indicated, e.g. to optimize fracture toughness. Work in this respect has mainlybeen concerned with the high-strength grades, since the 18Ni(200) grade exhibitsexcellent fracture toughness in the standard condition (see Table 6.1); very often, it wascarried out on plate or sheet products and combined with an evaluation of the effectof cold work. Figures 6.2 and 6.3 shov, the effects of cold work and maraging para-meters on the strength and fracture toughness of the !8Ni(250) and 18Ni(300) grades,respectively. The inverse relationship between these two properties is clearly apparent;il becomes even more evident when the Ku, values are plotted versus the yield strengths,as is shown in Figure 6.4 for the 18Ni(300) grade.

The conventional maraging grades do not respond to ausforming. Their response tomarstraining has already been dealt with in the preceding paragraph; a more completedescription of the effect of cold work on the mechanical properties of the 18Ni(25O)and l8Ni(300) grades is provided by Figure 6.5.

nJgOO 150D(12% OFFSET YIELD STRENGTH, WN/m? (200)

• •••: - 2000- • - •yV,- . -21DU.-":— : V 2200

WHi2300

120H

'••-.•:,-...-. , 7 7 1 .

^v^900»F(4BO°C|-2h

T . • : |2Q.°/.C.W.«S00°F(4B0°C

• • • • - : / • ] ' • :

— 40% C.W..900 °F(4B0'O3h

tgioo

7SO"F(400°C)-4h !

025% CW»900°Ft480°C)-9h

"î_JpiJ°F(«25 °C)-4hTO-1

70%C.W..900°FKB0°C)-3K!20

S00°n425tC)-30Dh

° |

B0 =

40260 270 1 b 0 29D 300 310 320

0 2 % OFFSET YIELD STRENGTH 103 psi

—'6075%CW..900°F|

340330

Fig. 6.4. — Strength/toughness relationsnip for 18Ni(300) gradesubjected to various combinations of cold work and ageing treatment.

Initial condition : • 0.06 in. sheet, 1500°F (8!5°C) - 15min, A.C. ] ., t j t u j :n a i) After H VV MAV-o 0.08, 0.12 or 0.24 in. plate cold rolled to L R ° J r ' a n d C C BUSCH' [6 '9]

0.06 in. sheet (25, 50, and 15% reductions) ) NOR ' J r " a n d t " t " l J'r 0.375 in. plate, t500°F (815°C) - 1 h, A.C. (long.). After E.P. GIIF.WICZ [6. IS].* 50 mi! sheet, 1500T (S15°C) - 1 h, A.C. (long.). After A.W. BRISBANE et at. \6.20].

71


300T LONGITUDUWL1 _ _ - . TRANSVERSE

LONGITUDINAL -

20 60 70 20 30%CQLD REDUCTION

50 60

2300

2200

2100

2000

Fi«. ft. 5. — Effect of cold work on smooth and notched properties of l8Ni(25O) and 18NU30O)grades \6.6}. initial condition (0.115 in. sheet) : lS00°F (815°C). Final ageing 900°F (48O°C) - 3 h.

6,1.2. High' and Low-Temperature Properties

The high-temperature tensile properties and impact strength of the I8Ni(250) andiNNii,300) grades are shown in Figure 6.6. The tensile properties of the 15Ni-9Co gradedeveloped for use at 10004F (540°C) have been included for comparison. The latter'sstress-rupture properties are given in Figure 6.7, together with three data points forshe ISNi(250) and lSNi(300) grades. The i5Ni-15Co grade recently developed for useat moderately high temperatures as a high-strength soft-magnetic material also exhibitsgood elevated-temperature properties combined with very high magnetic..induction [6.22].In particular, the stress corresponding to a creep rate of 1 % in 10,000 hours (as measuredfrom the slope of the creep curve at the 3500 hour point) at 850°F (455°C) is of theorder of ftO.OOO psi (420 MN/m-).

As regards cryogenic properties, the results of an investigation on I8Ni(200) plate areshown in Figure 6.8«. Although the A'u. value was considerably lower at —320°FI 19(vC) than at higher temperatures, this stea.l exhibited an excellent combinationof strength and toughness at all temperatures. The variations in the ultimate tensilestrength of l8Ni(?5O) and l8Ni(30O) sheet are shown n Figure 6.86. A nominal strengthof 450.000 psi <3i00 MN,m2) was obtained for the as-aged 18Ni(3OO) grade at —423°F( 253 C). but fracture toughness was correspondingly degraded. In the annealed, fullymartensitic condition, liotn steels retain very good toughness dovn to this temperature,while exhibiting strengths of 250,000 to 300.000 psi (1700 to 2100 MN/m*), depending>ui the grade. 1 his combination of properties is highly attractive for cryogenic appli-

t i ncations.

Bf-liarhur

Representative fatigue properties for the three grades, obtained from rotating beamtests oil ;uod bars* are given in Figure 6.9. The endurance limit increases with thetensile strength of Hie grede. Smooth ters of the !8Ni{200), 18Ni(250) and 18Ni(3OO)

6. Ni-Co-Mo MARAGING STEELS — THE CONVENTIONAL GRADES

T5ST TEMPERATURE,°C0 100 200 300 400 500

— 2000

Fig. 6.6. — Elevated-iemperature propertiesof 18NK250), 18Ni(300) and 15Ni-9Co (high-temperature) maraging grades. After [6.6] for250 end 300 grades, and S. FLOREEI andR.F. DECKER [6,21] for 15Ni-9Co grade.Condition: 1500°F (815CC)-1 h -f 9OOCF(48O:C) - 3 h (250 and 300 grade?); 1800sF (980°C),A.C. f 900°F (480'C) - 3 h (!5Ni-9Co frade).

Fig. 6.7. — Stress-rupture properties ofthe 15Ni-9Co high-temperature grade.After S. FLOREEN and R.F. DECKL'K [6.2/].Data for 18NK25O) and !8Ni(3OO) gradesincluêj for comparison are from [6.6].

200 400 600 eooTEST TEMPERATURE.°F

BOO 30 W 300RUPTURE LIFE, hours

230

TEST TEMPERATURE, °C-200 -150 -100 ~ -50

TEST TEMPERATURE,°C-200 -100

a e85"F(475"C)/8h -\o 900°F(480'C)/3h

60L

-ma -30B -200 -100 0TEST TEMPERATURE. °F

-320 -200 -WOTEST TEMPERATURE. °F

0 »70

a) 18Ni(200) plate. After C. VISHNEVSKY and E.A. STCI-OERWALD 16.23]. Condition: 1650°F(900°C) -2h. A.C.;1450cF ( 7 9 0 ° O - 2 h . A.C.; 900°F t480°C) - 2 h.

h) lSNi(25O) and (8Ni(300) shecl. After D.L.CORN [6.24]. Condition: 150DoF '815 C) - 1 h.A.C.; 90O°F (4S0cO - 3 h or 8 8 6 ^ (47«cO - 8 h.

Fig. 6.8. — Low-temperature properties of the 18Ni(200), 18Ni(250) and 18Ni(300) grades.

73


24U

20a

.160

n

in 80

40

(1 1 II i

ij.--—

1 I I I

e a ieNi(300)SO 18NK250)^ 18Ni(200)

i

I

1 1 l i 1

—

-

I 1 1

1600

1200

1s

SOU in" •

tn

ton

n101 n» n7

NUMBER OF CYCLES TO FAILURE

Fig. 6.9. — Typical rotating-beamfaligue properties of the l8Ni(200),18Ni(25O)and 18Ni(300)grades[6.5].Condilion: 1500°F(815°C)-1 h,A.C.

+ 900°F (480°C) - 3 h, A.C.

grades tested to 10s cycles have endurance limits of, respectively, 90,000 to 110,000 psi(620 to 760MN/m2), 90,000 to I i 5,000 psi (620 to 795 MN/m*) and 110,000 to130,000 psi (760 to 900 MN/mJ). The corresponding values for notched bars (A', = 2.2)are 40,000 to 50.000 psi (275 to 345 MN/m^), 40,000 to 55,000 psi (275 to 380MN/m2)and 40.000 to 60,000 psi (275 to 415 MN/ra2). Longitudinal specimens give values onthe high side of each of these ranges. As regards the fatigue crack propagation rate, ithas been shown that the 18Ni(250) and 18Ni(300) grades both obey the da/dN vs.(AA-'F-S relationship given in Chapter 4 (Section 4.2.3) [6.25].

Lowering the annealing temperature to 1400°F (760°C) increases the endurance limitof the three grades [6.26]. Similarly, nitriding 18Ni(250) parts has been shown toimprove their fatigue strength by some 20,000 psi (140 MN/m2) [6.27]. Finally, acomparison of the fatigue data for maraging steels and low-alloy steels of equivalenttensiie strength [6.28] has indicated that both groups are equally affected by the presenceor absence of compressive surface stresses, and that their endurance limits are compar-able, provided the parts are machined after ageing. Surface hardening by shot peeningalso greatly enhances the fatigue strength of the maraging grades [6.26, 6.28].


An exhaustive review paper on the stress-corrosion and hydrogen embrittlementbehaviour of maraging steels was prepared recently [6.29]. It is generally recognizedthat the 18";7Ni maraging steels compare favourably with other high-strength steels andoffer as good, or better, threshold plane-strain stress intensity values (KUcc) over awide range of strengths. In terms of critical crack size, maraging steels can toleratelarger flaws; conversely, for a particular flaw size, maraging steels are capable of with-standing greater loads without experiencing crack propagation.

Control of impurity elements is a significant factor in ensuring good stress-corrosionresistance. Marginal improvements were obtained by raising the carbon content of anl8Ni(300) grade from 0.03 to 0.06% [6.13]. However, using a statistical approach,carbon, and to a greater extent, sulphur, were shown to be detrimental to KXscc.Specifically, substantial improvements were obtained when the sulphur content wasdecreased to 30 ppm or less. Of the possible desulphurizing methods, electroslagremelting appears highly promising, sulphur contents of less than 10 ppm having beenobtained using high sulphur-capacity lime-fluorspar slags [6.30].

As regards the effect of heat treatment, a survey of available data on the effect ofannealing temperature has shown [6.29] that response to this parameter is not entirely

74

6. Ni-Co-Mo MARAG1NG STEELS — THE CONVENTIONAL GRADES

• 8 0 0 ^ 2 5 °G)/10hA 90Q?rTft80°C)/100ho 900°F(A86°e)/3V2h

100QI I 1 I

50 jf

ai

UJ20 c

in

10 <

2

15000 100TIME TO FAILURE, minutes

1000 10,000

Fig. 6.10. — Effect of ageing parameters on stress-corrosion susceptibility of 18Ni(3OO)grade in 3%NaCl at pH 6.3 and 1.7. After A.J. STAVROS and H.W. PAXTON [6.3!].

Initial condition : 1500°F (815°C) - 2 h, A.C.

consistent. Nor is the dependence of cracking resistance on grain size. However, itshould be noted that, in the case of the cast grade, replacing the 2100°F (1150°C)homogenization initially recommended for this steel by the lower-temperature homo-genization, overageing and annealing steps mentioned in Section 6.1 has resulted in amarked improvement of the stress-corrosion resistance in a 3.5%NaCI solution; thisimprovement has been attributed to grain refinement [6.3]. On the other hand, thereis general agreement that the standard ageing treatment produces the best stress-corrosion resistance. This is shown in Figure 6.10 for the 18Ni(300) grade in twodeaerated 3%NaCl solutions with different pH values; it is seen that time to failure isrelatively insensitive to stress intensity, but is highly dependent-on ageing parameters.The Klscc values for this steel were found to range from 10,000 to 15,000 psi y in.(11 to 16.5 MNirr3/2), independently of ageing treatment and environment (the twoNaCl solutions already mentioned, \N H2SO4, deaerated distilled water, oxygen-saturated 3%NaCl, and 3%NaC: + 1.5%Na2Cr2O7 at pH 6.1) [6.31]. The stress-corrosion resistance of the lower-strength maraging grades is appreciably higher; in3.5%NaCl solutions, a A'lscc value of 45,000 psi y in. (49 MNm-3/2) has been reportedfor the 18Ni(250) grade [6.32], while values as high as 100,000 psi V in. (HOMNm-"*)are associated with the 18Ni(200) grade [6.33]. Recent tests performed in 3.5%NaClon three experimental grades with increasing Ti contents have confirmed the dependenceof stress-corrosion susceptibility on yield strength : the steel with 0.65 %Ti (analogous to18Ni(300) : Y.S. 293,000 psi, i.e., 2010 MN/m2) had a KUcc value of 11,000 psi y >n-(12.1 MNm-3/2), whereas the steel with O.35Ti (analogous *o 18Ni(250) : Y.S. 254,000 psi,i.e., 1740 MN/m2) had a Kiscc of 47,000 psi v ' i n . (51.7 MNm-3/2) [6j4]_

In more practical terms, it has been stated [6.35] that the 18%Ni rnaraging steels canbe used at high stress levels in non-corrosive atmospheres (e.g., moist air) or at mediumstress levels in natural saline atmospheres (pH = 7.5); i I more acidic saline atmospheres(pH = 3.5), they should not be subjected to stresses of more than 60% of their yieldstrength. As regards the behaviour in solutions, these steels possess excellent stress-corrosion resistance, even at stresses close to their yield strength, in natural sea water(pH = 7.8); in deionized water, they can be used under stresses of up to 70% of their

75


yield strength. Finally, their stress-corrosion susceptibility at high stress levels in NaClsolutions saturated with H2S (/>H = 5) is large.

As regards the mechanism of stress-corrosion cracking in maraging steels, the literaturesupplies a considerable amount of indirect evidence favouring a hydrogen embrittlementmechanism rather than an active-path corrosion mechanism [6.29]. Concerning thistype of embriulement, many " direct" studies have been carried out on maraging steelseither precharged with hydrogen [e.g. 6.34] or tested in hydrogen gas. It has been shownthat, although these steels are subject to hydrogen embriulement. they tolerate greaterquantities of hydrogen than other high-strength steels for a given loss in ductility or agiven susceptibility to delayed fa lure; furthermore, they recover their original ductilitymuch faster on baking at 300T (150°C) [6.36].

6.3. Secondary Processing

Maraging steels are easy to cold work in the annealed condition, since they work hardenslowly [6.5, 6.6]. They can be reduced by substantial amounts before intermediateannealing is required; the reductions achieved are of 75% or more in cold rolling, ofup to 85 "a in wire drawing, and of 30 to 40% in deep drawing. The standard inter-mediate annealing temperature is 1500°F (815°C), with times of 1 hour per inch ofthickness; provided a 30% cold reduction is performed in the final operation, thistemperature can be raised to 1800°F (980°C). In addition to the three cold-workingoperations mentioned above, the maraging steels can be fabricated by processes suchas tube spinning, shear forming, explosive forming, hydroforming, bending and shearing.It should be noted that even though their work-hardening rate is low, the annealedductility as measured by elongation in 2 in. is comparatively low, ranging from about10 to 25"O. This is of particular importance in forming applications involving anappreciable amount of stretching; in the forming of deep-drawn shells or in cuppingoperations, drawing of the material rather than stretching should be the basis of tooldesign. However, stretchability has been found to be significantly improved by a prioroverageing treatment, and this has been put to use in commercial practice [6.37]. Finally,the combination of work-hardening rate and ductility of the maraging steels makesthem ideally suited to cold-heading operations.

Two recent review papers are available on the machining [6.38] and welding [6.39] ofmaraging steels. These steels are machined most easily in the solution-annealedcondition, in which the parts can be finished to their final dimensions since dimensionalchanges and distortion on ageing are minimal. Their machinability after ageing iscomparable to that of conventional steels of similar hardness. Grindability is essentiallythe same as that of ordinary constructional steels, provided a heavy-duty water-solublegrinding fluid is employed. The maraging steels can be torch cut; plasma arc cutting isa preferred method because of its efficient heat input [6.5].

One of the most outstanding features of inaraging steels is the ease with which they canbe welded without preheating in both the solution-annealed and in the fully heat-treated conditions. Only a post-weld ageing treatment is required to restore propertiesin the heat-affected zone and to develop good strength in the we •! metal. Gas-shieldedprocesses are generally favoured. TIG welding offers virtually no problems; on theother hand, MIG welding occasionally produces porosity, but this can be eliminate^by closer control of welding parameters or filler composition. Short-arc welding requiresthe use of pure helium shielding to produce welds with good mechanical propertieswhile maintaining adequate operability [6.5]. Details on filler metal compositions,processing parameters and other joining methods will be found in the review papermentioned above. Joint efficiencies of 90 to 95% are readily obtained with good weldductility and toughness.

76

7. Ni-Co-Mo MARAGING STEELS — THE ULTRA-HIGH STRENGTH GRADES

6.4. Applications

In keeping with the pattern followed in two recent articles [6.8. 6.40], the applicationsof the conventional maraging steels can be grouped under three headings :— aerospace, aeronautic and marine applications : these include rockel motor caseswith diameters ranging from 9 to 260 inches, and the load cells capable of measuringtheir thrust; pivots for the gimbal support of a missile trans-stage engine: the torsion-bursuspension system for the Lunar Rover Vehicle; flexible drive shafts for helicopters;aircraft landing-gear components; hinges for swing-wing planes; foil assemblies forhydrofoil ships; deep-submergence marine vehicles;

— structural and machine components : pressure vessels, components of liming mechanismin fuel injection pumps, index plates for machine tools; bolls and fasteners; barrels forrapid-firing guns; parts operating at cryogenic temperatures; springs and .estrainingdevices in miniature instrumentation;

— tooling applications : die-casting dies or components thereof; moulds for the plasticindustry; rams for extruding lead or lead-tin sheaths; cold-forming dies.

This great diversity of examples reflects both the excellent combination of propertiesand the very high reliability of the steels. It has been shown [6.41] that the reducedmanufacturing costs associated with their remarkable fabricability can more thancompensaie for the difference in price between the maraging steels and low-alloy steelsof equivalent strength. These factors should favour the development of new applicationsin all three of the above-mentioned fields.

7. Ni-Co-Mo MARAGING STEELS — THE ULTRA-HIGH STRENGTH GRADES

Although a 500,000 psi composition was delineated as early as 1966 (see Chapter 5,Section 5.1.2), the present chapter will be restricted to the 18Ni(350) and 13Ni(400)grades, since these are the only ultra-high-strength maraging steels on which work hasprogressed to the semi- or full-developmental stage.

7.1. Processing

Information on the preparation and processing of these two steels is scanty. A 5000 lb(2200 kg) heat of the 18Ni(35O) grade was satisfactory produced by vacuum inductionmelting and consumable-electrode vacuum remelting [/ / ] . The same process was usedsuccessfully to prepare a 2000 lb (900 kg) heat of the 13Ni(400) grade [7.2]. A laboratory-scale test on air- vs. vacuum melting was recently performed on the 18Ni(35O) grade; itwas shown that the air-melted heat was inferior in tensile strength and transverse ductility,but greatly superior in impact strength and longitudinal ductility [7.3]. This steel"scasting and hot-deformation behaviour is similar to that of the conventional maraginggrades, except that homogenization prior to forging should be carried out at 2250°F(123O°C) or below rather than at 2300T (1260°C), the homogenization temperaturecommonly used for the conventional maraging steels [7.1, 7.3].

The recommended heat treatment for the 18Ni(35O) grade is as follows : annealing for1 hour at 1500T (815°C) .ollowed by ageing at 900°F (480°C) for 12 hours or preferablyat 950°F (510°C) [7.1, 7.4] or 1000°F (540aC) [7.3] for 3 hours; lowering the annealingtemperature from 1500 to 1450°F (815 to 790°C) has also been recommended when priorforging is carried out at !900°F (1040C) instead of 2100°F (1 i50°C) [7.3]. The steel alsoshows an interesting response to ausageing [7.7]. As regards the 13Ni(400) grade, a recentinvestigation [7.5] has shown that homogenization at 1800°F (980°C) for 1 hour followed

77


by ageing for 4 hours at 980°F (525°C) appears to produce the best compromise betweenstrength and toughness.

Informavion on other aspects of processing is restricted to the 18Ni(350) grade. Coldrolling and wire drawing of this steel in the annealed condition was readily achieved,with reductions of 80% or more [7./]. Cold-rolled sheet was welded by the TIG process,using a filler wire of matching composition; sound, crack-free welds with good fracturetoughness were obtained, but the joint efficiency was rather lower than that achievedwith the conventional maraging grades [7./]. Machinability in the annealed conditionis excellent, and the machined parts are expected to undergo, upon full hardening, auniform linear contraction of no more than 0.1 %, with minimal distortion [7./]. Finally,successful preparation of the 18Ni(35O) grade by hot extrusion of pre-alloyed powderswas recently reported [7.6]; after heat treatment, the tensile strength of the powder-metallurgy material was found to be comparable to that of the forged alloy, while the

TABLE 7.1. — TYPICAL ROOM-TEMPERATURE PROPERTIES OF THE ULTRA-HIGH STRENGTH MARAGING GRADES

FormCondition

(all steps followedby air cooling)

U.T.S., 0.2 % Y.S., EL, R.A.,1 CVNimpact,

ft.lb /I03psi\/'in.MNm-'l'

N.T.S.

U.T.S.

I

Hard-ness,Re

Moduli,CN Ref.

0.5 in. plate (long.)

1 in. bar

1/4 in. bar ' vac. m.air melt

0.2 in. ho. .ig.rolled sheet | transv.

0.032 in. wire

Pr;alloyed pow-der extruded bar

'/2 in. bar

"/loin, plate

1500;F(815'O- 1 h- 950F(510°C) - 3 h.

150ODF(815eC)-l h- 900;F (48O"C) - 3 h.

DittoDitto

DittoDitto

as drawn"as drawn6

drawn11 and aged(900JF, 3 h)

2000^(1095=0-2 h-f- 950°F(510'C) - 3 h.

Hot rolled-f 900=F (480°C) - 4 h.

i !8OO=F(98O°C)-1 h- 900'F (480°C) - 4 h.

Hot rolled+ 900°F (480°C) - 4 h.

Hot rolled- l l i0 oF(6u0°C)-4h.

18O0°F(980°C)-lh+ 900°F (480°C) - 4 '„.

180O°F(980°C)-l h+ 980° F (525CC) - 4 h.

18Ni(350) grade

345 2379 331 2282 8.9 37.2

358 2468

329 2268

320 2206

332 2289

346 2386

224 1544

367 2530

429 2958

332 2289

352 2427

316 2179

308 2124

326 2248

335 2310

182 1255

287 1979

367 2530

320 2206

5

4.5

0.86

1.2

1.2

13Ni(400) grade

405 2792

390 26*9

368.7 2542

256.9 1771

350.6 H17

372.6 2569

395 2724

380 2620

362.3 2498

221 1526

350.6 2417

367.0 2530

5.2

4.3

0.63

1.8

43

24

38

16.5

25

27

1.5

5.5

0.7

6.4

7.2 9.8 —

12 16.3 —

T-- 9.5

10' 13.6

28 38 —

—

38 41.8

1 1 1 1

—

25 27.5

60.8 66.S

44.3 48.7

55.0 60.4

I0.58

ft = 7.5)

0.66K, > 12)

0.540.49

(.JCt = 18)

0.73(ATt=3.5)

—

—

—

—

—

58

—

61

59

62

51

63

64

•28.1 194(elasticity)10.7 74

.(rigidity)

30 207(elasticity)

[7.4]

17.1)

{7.4)

«

[7.1]

«

[7./]««

U.6]

[7.2]

«

[7.5]

» In two stages, with intermediate anneal at 1500°F (815°C).11 In two stages, with intermediate reversion treatment at U50°F (620°C). c Impact tested et —40°F (—40°C).

7. Ni-Co-Mo MARAGING STEELS - THE ULTRA-HIGH STRENGTH GRADES

impact toughness was considerably higher; this improvement was partly attributed lothe presence of continuous, Al-rich stringers in the material's structure.

7,2. Properties

7.2.1. Strength and Toughness

The room-temperature mechanical properties of both grades are listed in Table 7.1.This shows that the target of 350,000 psi is easily obtained through appropriate healtreatment; unfortunately, the associated N.T.S./U.T.S. ratio is well below unity, incontrast to the lower-strength maraging grades. The fracture toughness of this steelhas formed the subject of two major investigations [7.4 and 7.7]. It has been shownthat the Ku. value can be increased from about 36,000 to 40,000 psi v in. (40 to44MNm~3/2) by raising the ageing temperature from 950 to I050T (510 to 565"C).but this is insignificant when weighed against the attendant loss of strength [7.7].The work-hardening rate of the 18Ni(350) steel can be increased by ausageing or byapplying a reversion treatment [7./]. For instance, the yield strength of sheet treatedfor 1 hour at !500°F (815°C) + 24 hours at 1050°F (565°C) increases from 141,000 to277,000 psi (970 to 1910 MN/m2) on cold rolling to 25% reduction. Similarly, insertinga reversion treatment of 1 hour at 1150°F (620°C) between two cold-drawing operationsproduces wire with a tensile strength of 367,000 psi (2530 MN/m2), associated with anelongation of 1.2%; the excellent ductility is retained after ageing, while the strengthis raised further to 429,000 psi (2960 MN/m2).

The tensile properties and toughness of the 13NK400) grade, as listed in Table 7.1,show that annealing at 1800°F (980°C) slightly reduces the strength in the aged condition.The ductility of the semi-industrial heat [7.2] was quite good, while that of laboratoryheats [7.2, 7.5] was rather low. As stated above, the optimum compromise appears tobe obtained by homogenizing at \°r^p (980'C) and ageing for 4 hours at 9R0°F (525=C);under these conditions, the 13Ni(400) grade exhibits greater strength than the 18Ni(350)steel, combined with higher toughness. In contrast to the other maraging grades, thel3Ni(400) steel has a fairly large modulus of elasticity (30 x 10 psi, i.e., 207 GN/m2).which is fully retained up to 600°F (3I5°C) [7.2].

7.2.2. High- and Low-Temperature Properties

The effect of test temperature on the tensile properties and toughness of the 18Ni(35O)grade is shown in Figure 7.1. It is clear that this steel has a temperature capability of

TEST li'MPERATURE "C TEST TEMPERATURE, "C200 tOO

200 <00TEST tEMPERATURE^F

200 400 600 300TcST TEMPERATURE, 't

Fig. 7.1. — Temperature dependence of tensile properties and impact strength of 18Ni(35O) gradeAfter G.W. TUFFNELL and R.L. CAIRNS [7.1] (full lines) and A.A. IANNELLI [7.3] (dotted lines).

Condition Oh in. bar) : 1500oF (815"C) - 1 h, A.C. + 900°F (480°C) - 3 h, A.C.

79


HQLDIN'3 A \ G TEST TEMPERATURE.'C

200 h HOLDING0 200 «ffl 600 600 OOO

HOLDING A N 0 TEST TEMPERATURE,*F

n ioo BOOTE5T DURATION, minutes

fill

LABORATORY AlR

800 'F/Bh *3 T a

3.5Y. NaCI SOLUTION

BOO'F/Bh O • ®

900'F/Bh a A »

95O'F/3h • • O

Fig. ~.2. — Effect of test temperature, with or with-out prior holding, on tensile properties of 13NK4O0)crade. After W J . BotscH and T.W. COWAN [7.2).Condition O!z in. ban : hot rolled - 900"F - 4 h. A.C.

Fig. 7.3. — Effect of ageing treatment on sub-critical crack growthresistance of 18Ni(350) steel in laboratory air and 3.5% NaCI solution.

After C.S. CARTER [7.7].Initial condition (double anneal) : !70(TF m S - C ) - I n - ' - 1 W F ( 8 I 5 C ) - I h.

about 900;F (480 C). the drastic loss of strength experienced at !000°F (540°C) beingdue to overageing and reversion to aastenite. The constancy of the hardness, tensileelongation and impact strength up to 900°F are worth noting. In contrast, the impactresistance of fatigue pre-cracked Charpy V-notch specimens (not shown in the figure)rises considerably between R.T. and 5003F (260JC); this reflects an increased resistanceto crack propagation, which would be an advantageous factor in tooling and dieapplications at moderately high temperatures [7./].

As regards the l3Ni(400) grade, a study aimed at determining its maximum servicetemperature [7.2] gave the results summarized in Figure 7.2. The temperature capabilityin tnis case lies near 800=F (425'C). Prior holding at the test temperature for 200 hoursdecreases the slope of the strength vs. temperature curve; the resulting gain in strengthis accompanied by only a slight decrease in ductility. A similar effect was observedin the case of the 18Ni(350) grade [7.8].

7.2.3. Other Properties

After annealing at 1475T (800°C) and ageing for 10 hours at 900°F (480°C), the18Ni(35O) grade has a 10» cycle fatigue limit of 110,000 psi (760 MN/m^) [7.8]. The high-cycle fatigue strength of this steel can be improved by ausforming and marforming [7.9].

The stress-corrosion resistance of the same steel in a 3.5%NaCI solution was determinedfor three ageing treatments (Fig. 7.3). The results showed that the material aged below900°F (480'C) was very susceptible to sub-critical crack growth; in fact, prematurebrittle fracture occurred in unnotched tension specimens loaded at a slow strain ratein laboratory air. Such sub-critical crack growth was attributed to a hydrogen embrittle-ment mechanism. Ageing at 900 and 950"F (480 and 510°C) improved the thresholdstress intensity by a factor of 2.

7.3. Applications

The 18Ni(3:;0) grade has already found applications as cold-heading dies, cam followers,forging dies, extrusion rams and core pins in die-casting dies [7.8]. As regards the13Ni(400) grade, the combinations of strength/stiffness and strength/toughness/form-ability suggest its possible use as tooling fixtures and fasteners, respectively [7.2]. Finally,the hardness and elevated-temperature capability of bath-steels should favour their usein the field of bearings. -"

80

8. STAINLESS MARAGING STEELS — PHYSICAL METALLURGY

8. STAINLESS MARAGING STEELS — PHYSICAL METALLURGY

8.1. Background

Cobalt-containing stainless maraging steels were developed at the incentive of users invarious fields — machine construction, aerospace and chemical industries, naval engineering,etc. — who required steels possessing good resistance to corrosion or tarnishing, associatedwith strengths in excess of about 140,000 psi (1000 MN/m-), and preferably approachingthose of the Ni-Co-Mo maraging steels. Before dealing with their physical metallurgy,it seems appropriat' to define.the position of these steels with respect to other stainlesssystems.

The basic prerequisites for the cttainment of a high-strength stainless steel are that itshould contain sufficient chromium and exhibit a martensitic structure. As regards thechromium content, it is well established that this must be greater than 10%, or possibly12%, if adequate corrosion resistance is to be achieved. However, an upper limit ofapproximately 17%Cr is imposed, mainly by the second requirement, viz. the feasibilityof obtaining an essentially martensitic structure after heat or thermomechanical treatment.This condition also implies that the composition of the steel be balanced so as to ensurea microstructure essentially free from 8-ferrite at the austenizing temperature, and anMs point preferably above room temperature. As stated in Chapter 2 (Section 2.3.1),cobalt is an efficient austenite stabilizer in 12%Cr steels and can thus be used to controlthe S-ferrite content; in addition, cobalt only slightly depresses Ms and is therefore lesslikely than other y stabilizers to cause an excessive increase in the amount of retainedaustenite. In addition to the strengthening provided by the martensitic transformation,high-strength stainless steels are usually hardened through precipitation of carbides and/orintermetallic compounds, which occurs during the so-called " fifth " stage of tempering(c/. Chapter 2, Fig. 2.16).

Straight chromium martensitic stainless steels (with moderate strength) are exemplifiedby A1S1 types 420 and 431 (Table 8.1). In the former, the martensitic structure is ensured

TABLE 8.1. — NOMINAL COMPOSITIONS Or SOME COBALT-FREE MARTENSITICAND CONTROLLED-TRANS.FORMATION STAINLESS STEELS (in wt.%, bal. Fe)

Designation

AISI420AISI431

AISI 616 (422)AISI61917-4PHAM-363IN-736

PH 13-8 MoCustom 450Custom 455

PH 15-7 Mo- 17-7 PH

AM-350AM-355

PH 14-8 Mo

C

0.15 min.0.2 max.0.230.30.040.05 max.0.010.040.05 max.0.03 max.

0.070.07o.o;0.130.04

Cr Ni Mo Cu

Marteiisitic stainless

1316

1211.41611to12.S15.011.5

20.80.34.34

108.16.58.5

—12.8

.22.20.8—

—

—3.3

1.52.2

Tj

steels

—

—

10.2

1.2

Al

—

—

0.31.1——

Controlled-transfonnation steels

15.11716.515.514.4

7.17.14.34.38.2

2.2

2.82.82.2

1.11.1

—i.l

Others

—1W, 0.25V

0.25V0.25Nb

_

0.7NbO.25Nb

0.1N0.1N

Owner of trade name

———

Armco SteelAllegheny LudlumInternational NickelArmco SteelCarTechCarTech

Armco SteelArmco SteelAllegheny LudlumAllegheny LudlumArmco Steel

81

COBALr-lONTAlNlNCi HIGH-STRENGTH STEELS

by a relatively high carbon content; in the latter, which has a higher chromium content,nickel is used together with carbon to control the microstructure. In both steels, part ofthe strength is derived from the precipitation of chromium carbides, and yield strengthsas hi eh as 220.000 psi (1500 MN/m-) may be achieved depending on the heat treatment.Many l2",,Cr steels hardened by means of alloy carbides (through addition of strongcarbide-forming elements such as tungsten, molybdenum and vanadium) have beendeveloped to meet the current need for stainless steels with good properties at temperaturesup to 1200 F (650 C) [8.1]. Although these steels can exhibit strengths of the order of200.000 psi (1400 MN m-k they are most often treated to relatively low room-temperaturestrength levels {ri:. yield strengths between 120.000 and 130,000 psi or 840 and 910 MN/m-),with a vietv to optimizing their properties at higher temperatures. Typical examples ofthis class are AISI 616 and 619. Several 12%Cr steels which derive their strength from theprecipitation of intermetallie compounds have also been developed; as shown in Table 8.1,they are low iii carbon, but contain elements such as molybdenum, niobium, copper,titanium or aluminium. As an example, 17-4 PH can be heat treated to a room-temperature yield strength of 170,000 to 200,000 psi (1200 to 1400 MN/m'-).The so-called semi-austenitic or controlled-transformation steels constitute yet anotherclass of high-strength stainless steels. These steels, the nominal compositions of whichare also listed in Table 8.1, have an austenitic microsti ncture when in the anneatedcondition, but transform to martensite after ageing, refrigeration, or deformation. Sub-sequent hardening takes place by precipitation of intermetallic compounds and carbides,giving room-temperature yield strengths up to about 200,000 psi (1400 MN/m-),

Although cobalt has been shown to be a useful alloying element in both straight-chromium and controlled-transformation steels {if. Chapter 2, Sections 2.1 and 2.3.7respectively), no cobalt-containing versions of such steels have been commercialized. Onthe other hand, quite a number of cobalt-containing martensitic stainless steels hardenedby means of intermetallic compounds, and possibly of alloy carbides, hav? been developedover the past few years. It will be seen (Table 8.2) that all these steels contain chromium(10-16%), cobalt (5-20%) and molybdenum (2-5.5%), and that a number also containnickel (I to 8°/J, for reasons that will be discussed later on. AFC-77, its Nb- and Ni-containing modification (" Alloy B "), and Pyromet X-12 contain approximately 0.15%Cand are thus receptive to both carbide and intermetallic-compound strengthening. All theother steels have very low carbon contents and thus depend essentially on the precipitationof intermetallic compounds for their strengthening; the predominant tendency towards

TABLE 8.2. — NOMINAL COMPOSITIONS OF COBALT-CONTATNING MARTENSITIC STAINLESS STEELS (wt.%, bal. Fe)

Designation *

AFC-77AFC-260 (C50)

" Alloy B "AM-367

D. 70Pyromet X-12Pyromet X-15Pyromet X-23Ultrafori 401Ultrafort 402Ullrafort 403

C

0.150.080.160.03 max.0.03 max.0.120.01 max.0.03 max.0.02 max.0.02 max.0.02 max.

Cr

14.5!5.5141412

10.5i ;10

1212.511

Ni

2.01.03.54.3

——7.0

8.27.6

C o

13.51313.515.514.5

62010

5.35.49

Mo

5.04.35.02.04.0

4.82.95.5

7.04.24.5

Cu

—

—

i 3

—

—

Ti

—

0.5

N b

0.140.22

present

—

0.80.50.4

—

—

Others

0.5V—

Al, B, Zr

0.08N

—

B, Zr0.05A10.15A1

Owner of trade name

CrucibleCrucible

Allegheny LudlumEnglish Steel

. CarTechCarTechCarTech

DEWDEWD E W

Yearannounced

19631967197119631964

196119671973

196919691971

Some of these steels, suci: as Pyromet X-12 and Ultrafort 402, have been superseded by later grades.

82

8. STAINLESS MARAG1NG STEELS — PHYSICAL METALLURGY

10 0 2 4 S B 10 0 2 4 6 10 0 ' 2 i E 8 10

Fig. 8.f. — Influence of cobalt on the constitution of Fe-13Cr andFe-17Cr alloys at two carbon levels. After D. COUTSOURAOIS [8.3].

such low carbon contents is related to the wish to obtain as tough a martensite as possibleby favouring the formation of a low-carbon lath martensite rather than a high-carbontwinned one. Although, strictly speaking, the terminology " stainless maraging steel "should be restricted to the low-carbon grades, it is generally considered as applying tothis family of high-strength steels as a whole, and will be used in this general sensethroughout this and the following chapter.

Finally, mention should be made of the existence of a number of heat-resisting steels withcobalt and chromium contents comparable to those shown in Table 8.2 [S.I]. However,they will not be considered explicitly here, since the emphasis in the present volume is onroom-temperature properties.

8.2. Effect of Alloying Elements on Equilibrium Structures

Annealing of steels containing from 10 to 15%Cr will produce, at temperature, a fullyaustenitic (y), a mixed austenitic-ferritic (y + a) or a fully ferritic (a) structure, dependingon the annealing temperature and the alloying element (Cr, Ni, C, Co, etc.) content.When the steel has been at least partly austenized, the microstructure after cooling willdepend essentially on the cooling rate and the alloy element contents. Since the structuralrequirements for high-strength stainless steels impose that a virtually ferrite-free structurebe attained on austenizing, it is appropriate to review here the effect of various alloyingelements on the ferrite/austenite stability relationship at high temperatures.

The ability of cobalt to stabilize the aiistenite in Fe-Cr alloys was recognized as early as1928 [8.2], but was only evaluated quantitatively at a much later date [S3]. Figure 8.1shows the effect of adding up to .10%Co on the equilibrium structure of Fe-i3Cr andFe-17Cr alloys containing 0.05 and 0.2%C. It is evident that addition of relatively smallamounts of cobalt to 13%Cr steels broadens the y field without greatly depressing thea ;=i v or a ^ a + y equilibrium transformation temperatures. The latter point is ofparticular importance in steels intended for high-temperature applications in whichtransformation to austenite during service is undesirable.

Figure 8.1 also shows that low-carbon 13%Cr steels are likely to form ferrite at the highestaustenizing temperatures. This constituent, which is normally designated S-ferrite, isdesirable in restricted amounts in controlled-transformation stainless steels, but must beavoided in straight martensitic ones because of its deleterious effect on strength. Forthis reason, a number of studies have been performed in order to determine quantitativelythe influence of various elements on the S-ferrite content of stainless steels. Figure 8.2

83

I u H . U T - l O \ l \ l N t \ C , UK.(I

CARBON COMTEMT. %0 3 Oi 05 06 07 08

ALLOYING ELEMENT CONTENT, %10' O

LOG TIME (seconds)

1 in - ;. 1 lieu of \,in.'U> .ill.'wni: clc-mcni> i>n the •S-ferrue comcnt of a O. IC-t"Cr steel. After K.J. lRMxt t-( ,i/. [S.4].

1 iiz K..I, isiMhermal trans!t>rma!ion of ^ ferrite inlou-carbon Al-C-77 and in AKC-2M) after austeni/ation for1 hour a! 2:00 F t i ; 0 5 C ) . After D. WF.BSTIR \S.!U\.

illustrates the etfect of se\eral v-stabili/inj elements on i.hc -vfcrrite content of a 0. IC-l7Crsteel. According to this figure, cobalt is less effecti\e in suppressing ^-ferrite than eithercarbon or nickel, but is more so than copper or manganese. Similar data have beenestablished for 0.lC-!"Cr-4Ni steels [X.-f]: the results can be summarized as follows :E l c m c m U " . , . eM-eru C and N . 0.1", , i N C Ni t o Cu Mn W Si Mo Cr V Al

Change in ^-ferriie contentl i n a b s o l u t e " „ ) - 2 0 - I S — i O - h — 3 - I - 8 - 8 I I - 1 5 • 1 9 - 3 8

In this case, cobalt appears to be 0.6 times as effective as nickel in suppressing >vferrite.*The same ratio was derived from a stud\ of the efff-ct of cobalt on the o-ferrite contentof O.lC-lKCr steels [SJ]. In fair agreement with the above data, the following relationshipwas established for the equivalent nickel and chromium contents of steels containingup to 12";lCr [8.5] :

Ni,., - -30r , ,C • ",,N) • O.SC'.Mnl - (",,Nh -0.7(",,Co) (K.I)Cr,,, = (",,Cr) - 1.5(%Si) - (°,,Mo) -- l.5(%Ti) (8.2(

Other investigators [8.6] arrived at the same conclusion, except that they assigned highervalues to the coefficients of cobalt (0.8) and titanium (2.0) in the respective equations

In commercial steels, S-ferrite may form on heat treating at elevated temperatures,depending on the balance of austenite- and ferrite-stabilizing elements. For instance, ithas been shown that a standard heat of AFC-77 austenizeu at 1900"F (1040 C) for 1 hour,oil quenched, and refrigerated for 'K hour at —100 F (—73 C) contains 5% o-ferrite[8.7]; higher austen;zing temperatures produce increasing amounts of 8-ferrite, e.g. 15%after a 2400"JF (1315 C) treatment [«.#].

Since high austenizing temperatures are unavoidable in the case of complex compositionssuch as AFC-77, because of the increasing amounts of hardening elements that have togo into solution, special treatments have been developed to decrease the ^-ferrite contentwithout lowering the austenizing temperature. One such treatment consists in temperingfor 2 hours at 1400 F (760 C) prior to austenizing at 1900 F (1040 C); this reduces theS-ferrite content from 5 to 2",, [8.7]; the reason for the effectiveness of the preliminarytempering step is related to the removal of alloying elements from the high-temperatureferrite phase through precipitation of carbides and intermetallic compounds. Anotherpossibility is to austenize the steel at the required temperature (2100 to 2200°F, i.e., H50

• Note added at proof ,tage. The data in tilts section and in Section 8.3.1 were confirmed by a recent investig-ation [HULL. F.C., WelJina J.. Res. .Simp/.. 52, I93s <I973>1 on cast complex 12-24",, Cr steels, in which cobaltwas found to be 0.5% as effective as nickel in suppressing S-ferrite formation, and 0.25 times as effective as nickelin depressing M*.

84

K. S T A I N U S S M A R A G I N G S U I I S P H Y S I C A L M E T A L L U R G Y

to 1205 C) and then to soak it for 1 hour at about 1900 F (i040 Cl before cooling to roomtemperature, in order to give the residual ^-ferrite sufficient time to transform to austenite[iS.V|. The •S-ferrile present after high-temperature austenizing effectively pins the austenitegrain boundaries, resuming in a line austenite grain size: after isothermal removal of theferrite. a marked increase in austenite grain size is observed [H.H!]. When performedcorrectly, this two-stage treatment not only eliminates ^-ferrite. but also results in a bettercombination of strength and toughness through refinement of the austenite grain (seeSection X.fc.4).

The kinetics of the ^-fertile to austenite transformation were determined on a lower-carbon modification of AFC-77 (0.1 I ",,C instead of 0.15",,), as well as on AFC-260 [8.10].The TTT curves for this transformation are shown in Figure 8.3; both steels were austenizedat 2200 F (1205 C) for 1 hour and then isothermally treated for increasing times attemperatures between 2200 and 1600 F (1205 and 870 C). The transformation curvesfor AFC-77 with the usual carbon content would be shifted to the left of those shownin the figure.

8.3. Transformation Temperatures and Structures

Since high-strength stainless steels are designed so as to achieve a fully niartensitic structureon cooling to room temperature, the effect of composition on the A/.< temperature is ofprimary importance (a general discussion of this topic will be found in Chapter 2.Section 2 3.1). Similarly, the retained austenile content, as controlled by the position ofthe Mf point, can markedly affect the properties obtained. Finally, the transformationof martensite to austenite on tempering or ageing (or during service in the case of steelsused at high temperatures) is also of great significance. The present section is devoted to adiscussion of these various aspects.

8.3.1. Manensilk Transformation

The low carbon contents of the stainless maraging steels under review result in theformation of lath martensite, the occurrence of twinned martensite being ratherunusual. The significance of achieving a lath martensitic structure, particularly withrespect to precipitation kinetics and precipitate dispersion, on the one hand, and totoughness, on the other, was emphasized in Chapter 2 (Sections 2.3.3 and 2.3.6), andneed not be reconsidered here. As a matter of fact, all the steels listed in Table 8.2 exhibita lath martensitic structure (of the type illustrated in Chapter 2. Fig. 2.10) on quenchingfollowing austenization. However, in the higher-carbon steels such as AFC-77, limitedamounts of twinned martensite are often co-present (Fig. 8.4).

Fig. 8.4. — Thin-foil iransmission electron micrographof AFC-77 after quenching from 20!0°F (lI003C)and refrigeration at —lOO^F (—73°C) for 15 hours.After E. DIDERRICH et at. [8.11]. ••: 20,000

85

; \l > N I \ l \ I N i • I I H i H N | R | N V , I I I

I it:. S*. - V.lYivl of lanous alln>inji clement-;.•n the \ / , leniperaiure of :i U.i)3C-l-C"r Mcd

Ada I ' M I U M M O N U [,S'..5].

Ha.se oiniposiiion

(IO.1C.-4Nll>.03L"-i:i r0.0.iL~-IXr-4Ni( IDH' - I ' l r-4Ni-15t iOO.K"-i:t'r-4Ni-l5L'i

I lemeiil K.mj:e

I ' r (1- \2 » l . " , ,S i 4 - SI'M 0- I*

Mi- l>-f

1 I 0-1.1

"C 2 4 S" 9 10 1? ' i <6ALLOYING ELEMENT CONTENT. * t •/.

As regards the " martensitic " transformation temperature, there is some evidence that,in the t\pes of steel considered here, it i* dependen on cooling rale. As an example, theV/r temperature of a l4Cr-20Co-5Cu steel was fuLiid tn he lowered from S80 F (470 C)tn 565 F (2^5 Cl when the cooling rate was raised from 720 F min (400 C min) to 90 F seci 50 C sec) [S.I2]. In the subsequent discussion, the M, temperature should be understoodas corresponding to cooling rates above 90 F sec.

The effect of chromium, nickel, cobalt, molybdenum and titanium on the Mn temperatureof l2",,Cr steels was determined for a series of 0.03 ",,C compositions solution treated toobtain a completely austenitic structure. The results are given in Figure 8.5 togetherwith the ranges of contents investigated. Using these data, the following relationshipbetween M, and composition was established [8.5] :

M,( F) = 1530 -52( o , ,Cr)-70(" , INi) — 9(";,Co)~65(%Mo) — 0(°nTi),V/.,( C) = 832 - 29("oCr) — 39(",;Ni)- -5(%Co) — 36(%Mo) — 0(%Ti)

An interaction was noted between molybdenum and titanium. When .idded singly, thelatter has a coefficient of 0: in the presence of molybdenum its coefficient i?. —100(F) or—55( C). According to Eqn. 8.3, cobalt is 0.13 times as active as nickel in depressing M8.

A similar equation has been established, by regression analysis, between Af50 (thetemperature at which 50% austenite has transformed) and composition, for steels contain-ing 0.01 to 0.03°,,C, 9.2 to l3.4",,Cr. 5 to 7.7';,,Ni, 5.6 to 12%Co, 2 to 4.8"/oMo and 0.1to 0.7 %Ti [8.6] :

M50CF) = 2 8 4 2 - 8 5 ( % C r ) - 157(%Ni)-50(%Co)-65(%Mo) + 205(%Ti)M50(°C) = 1561 —47(%Cr) —87(%Ni) —28(%Co) —36(%Mo) + 114(%

In this case, cobalt appears to be 0.32 times as effective as nickel in depressing the A/;o

transformation temperature. The large positive titanium coefficient is possibly a resultof the interaction of this element with carbon and nitrogen.

86

S S I M M l . s s M M . ' \ ( , l \ ( , S i l l I s \ l M I I M l U R < A

Both Fqns. S.3. and H.4 can be converted to so-called •• Austenite Retention Index "'(A.R.I.I relationships \N.I3] which embody weighted factors for each element, based onits '.-fleet on M,, as compared with that of nickel for which a factor of I is arbitrarilyassigned. Rewriting them in this way gives, respectively :

A.R.I. Ni • 0.74 Cr • 0.13 Co • 0.42 Mo (8.5)A.R.I. Ni • 0.54 Cr • 0.32 Co • 0.41 Mo <H.6>

In this form the> can be compared with tne original A.R.I, equation used lo developi'yromet X-23 [X.U] :

A.R.I. --= Ni • 0.8 Cr •• 0.3 Co - 0.6 Mo (S.7)

It is seen that agreement between these three relationships is reasonably good, considerinethe approximations inherent in this kind of treatment.

In conclusion, it clearly appears that the ell'ecl of an individual addition element on theMt temperature can vary quite extensively with the overall composilion of the steels. Thisis illustrated in Table 8.3 for six major alloying elements (graphical representations of theinfluence of molybdenum and titanium will be found in Chapter 5. Fig. 5.5). It is seenthat, in stainless maraging steels, the most probable changes in M.< are —72 F ( -40 C) per" I . " , , Ni. —45 F (—25 Cl per vvt.",, Cr. —9 to —18 F ('—5 to —10 C) per wt.",, Co,

and — 72 F ( 40 C) per wt.",, Mo. As already stated, the effect of titanium appears tobe highly dependent on interaction phenomena. Finally, copper seems to raise theA/,,, temperature, at least in !4Cr-20Co steels.

TABLE 8.3. — EFFECT OF VARIOUS ALLOYING ELEMENTSON M, TEMPERATURE OF STAINLESS STEELS

Alloyingelement

Ni

C.

Co

Mo

Ti

Cu

Steel typesor compositions

Standard stainless0.IC-18O

12Cr~I0Cr

5-10Cr, 5-8.5Ni5-IOCr, 8.5-11.5Ni

Standard stainless4Ni

Co-Ni6"Ni

0.1C-18Cr12Cr-4NilOCr-Ni

12Cr-4Ni-15ColOCr-Co-Ni

l4Cr-20Co-2Cu

12Cr-4Ni-15Co(OCr-Co-Ni

14Cr-20Co

EITect on A/, per \ u . ° , ,

F C

— 110 —61—72 —1()—70 —39

—I5f>* —87*—72 — / 0—36 —20

—76 —t:—52 — 29

—85* —47*—41 —23

— 17 —9.5—9 —5—SO* —2S*

—65 —.«—65* — 36*—S1 —J5

0 to —99 0 lo —5.5

+205* -114*

+ 13 +7

ReJerence

[S.I-4]

18.3][S.S]

l«.6Jin-15]IS-15]

IS. 14][S.S][S.6]

[S.I 5]

[SJ]

[H-5]IS.6]

18.5][S.6]

[S.12\

[8.5][8-6]

[S.I 2]

Effect on M;Q temperature.

87

I OH U I -l O \ I \ 1 M \ C , H K . I I S I K I M i U I SI M ' l S

m

1 10 100 1000

HOLDING TIME AT INDICATED TEMPERATURES, minutes

Fie. S.h. IsD'^frrnal austenite reversion in C'r-C'o and Cr-Ni steel*.After K. B I S C M I D T et til. \H.17].

Initial condition : 142(1 V 11 .i5() Cl - I h, A.C.

20" Z1 22 "23A.R.I (Ni-0.aCr»a.3Cr.0.6Mo)

J10002t

Fig. S.7. — tiffed of A.R.I, on austenile content andultimate tensile strength of Fc-Cr-Co-Ni-Mo steels.Alter R.L. CATON and G.N. MANIAR [8.I3\. Condition:1700 V (925 O - Ih. VV.Q. • 1000 F (540 Cl-4h, A.C.

8.3.2. Austenile Reversion

The transformation of martensiu to austenite during exposure at elevated temperaturemay be a serious limitation in thi- choice of tempering, ageing, or service temperatures.For example, the As temperature of a O.IC-l2%Cr steel is approximately 1365°F (740°C),so that tempering temperatures up to 1300 F (700=C) can be used without any adverseeffects. However, if appreciable quantities of alloying elements such as nickel are added,the As temperature may be depressed to well below 12003F (650°C), leading to considerabledifficulties during heat treatment.

The effect of various elements on the As temperature of the 0.IC-12%Cr steel referredto above is as follows [S.I6] :

Element (1%) Ni Mn Co Si Mo Al V

Change in A,, G h C — 54, — 30 —45,-25 — 9 , - 5 ^-36,-20 -45 , +25 +54, +30 +90, +50

As already stated, cobalt only slightly depresses As, so that the martensite in chromiumsteels having compositions balanced essentially with cobalt can be expected to exhibitconsiderable stability with respect to reversion. This is exemplified in Figure 8.6, in whichthe behaviours of Fe-19Cr-15Co and Fe-14Cr-8Ni martensites on holding at varioustemperatures are compared. The curves clearly indicate that, at a given temperature,austenite reversion has proceeded much further in the Fe-Cr-Ni than in the Fe-Cr-Costeel; conversely, higher temperatures are required to form a given amount of austenitewithin a specified time i,. the Fe-Cr-Co martensite. In both cases the time dependenceof the reaction is evident. The stability of martensite in cobalt-containing stainless steelsis also demonstrated by the fact that no austenite reversion occurs in Ultrafort 401 afterexposure for 1000 hours at 800°F (425°C), and little reversion is observed after holdingfor 100 hours at 930°F (500°C) [8.6].

88

S I . M M I S S M A K A d l N t i S I L I . l . S P i n S l ( A t M f T A I L . L ' R C , Y

The stability of the reverted ausienile on subsequent cooling is strongly dependent onco-..position. For instance, it has been shown [8.17] that the ausienite formed in theabove-mentioned Fe-iyCr-15Co steel during exposure for up to 1000 hours in the 8-*0-1290 \r (450-700 C) temperature range completely transforms to martensite on subsequentcooling; on the other hand, the reverted ausienite formed in the Fe-l4Cr-KNi sieel appearsto be quite stable. It should be staled, however, that this difference in behaviour is partlydue to the increase in the Mf temperature of the former as a result of the more extensiveprecipitation reactions which it undergoes during the isothermal treatment.

8.3.3. Retained Austenhe

In controlled-transformation stainless steels, up to about 10",, retained austenite can betolerated before any appreciable decrease in yield strength is observed [X.!8]. Amountsof similar magnitude are often accepted when working out the compositions of stainlessmaraging steels, since this enables the quantities of alloying elements added lobe increased.On the other hand, the retained austenite content is an important factor in obtaining highfracture toughness in martensitic steels [8.B. 8.19]. In the Fe-Cr-Co-Ni-Mo system, theamount of austenite that ensures optimum fracture toughness without impairing yieldstrength is comprised between 5 and 15",,; this amount normally includes both retainedand reverted austenite, as produced respectively on cooling after solutioning and onageing [8.13].

The principal parameters controlling the retained austenite content are composition andheat treatment. A quantitative assessment of the effect ot various addition elements onthe retention of austenile is provided by the austenite retention index discussed inSection 8.3.1. The relationship between the A.R.I, calculated on the basis of Cqn. 8.7 andthe retained austenite content in Fe-Cr-Co-Ni-Mo steels is shown in Figure 8.7. Therather large scatter in the data points has been attributed to the fact that a linear equationwas used to describe a non-linear phenomenon [8.13]. This figure also shows that theoptimum strength of these steels corresponds to a retained austenite content of about 10",,.

The effect of the austenizing temperature on 'he retained austenite content in AFC-77is shown in Figure 8.8. It is seen that the amount of retained austenite increases withincreasing austenizing temperature. The reason for this is that high austenizing temp-eratures increase the alloy content of the matrix by dissolving alloy carbides and inter-metallic particles, thereby depressing Ms and My.

30

1700

950AUSTENIZING TEMPERATURE,'C_BOO 1050 1150

1800 1900AUSTENIZING TEMPERATURE, ' F

2000 2100^

Fig. 8.8. — Effect of austenizing temperature on retained austenite content in AFC-77 after lemperingat indicated te.nperatures. After D . WEBSTER [8.19]; insert after E.J. DULIS and L. HABRAKEN [8.8].

Condition : 7aust.- I h, O.Q. + refrigeration + ?~icmp.- 2 + 2 h.

89

I \ l 1 • M \ 1 M N I i U l i .11 M K! Ni . 1 II S.i I I I s

of.i'rct =Jat aetorrr- twjn in (emperwj conditionCondition 14GQpn753'CI-t.h»de!cyrr,a-

I :- \ l < '-" WlllMIB I'1*..'")

I he elVcel nf tempering ieiv.perature on the retained auMenile content m AI C - 7 7 is alsosho'.vn in Fiiiunj VS. The insert m this tigure i i ius the results of another investigationwhich involved a sinelc auMeni/ing temperature but a broader range of temperingtemperature-. Increasing tempering temperatures result in decreasing retained austemtecontents, due to she precipitation ol larger amounts of carbides and mtermcialliccompounds, which raise- the \t. tempera ure. If. prior to tempering, the steel is eitherausformed 5(1",, at IOC" 1 (540 Cl or cold worked H>",,. the retained ausiemte contentsafter the double tempering step arc mdepcndenl of tempering temperature up to 1011(1 I'(540 Ci . and equal to 15 and 5",,. respecti\el\ |V /v j .

S.4. Gra'.n Size

The favourable efleet of decreasing grain size on (he yield stienglh and fracture loughnessof high-strength steels prompted the initiation of several studies aimed at controlling thegrain size, particularly in AFC-77. Various sequences of therinomechanical trea'mentsused to obtain grain refinement in ihi.- sieel are shown in Figure S.y a. Sequence A consistsin deforming ihe austenite by at least 70",, in the 2OOO-I5OO F (II00-SI5C) temperaturerange, cooiing to room temperature, and annealing at 1800 F (980 C). Using the samedeformation conditions, efficient grain refinement is obtained when the stee! is reheatedfor 30 minutes a; 1600 to 1800 F (870 to 980 C) immediately after reduction (sequence B).Ausforming at 1000 F (540 C) to 50",, reduction prior to annealing at 1800 F (sequence C)gives a duplex mierostruciure. Grain refinement is also obtained through room-temperature deformation (up to 35",,) of the as-quenched or tempered martensite(Fig. 8.9 h); deformation of tempered martensite (sequence F) is obviously more effective.If deformation of the tempered martensite is carried out at 1400 F (760 C). the grainafter final annealing is much larger (Fig. 8.9 c).

The above austenite grain refinement is related to the formation of deformation voids atthe interfaces between matrix and hard carbide particles. The voids remain stable up tofairly high temperature (1875 F, i.e. 1025 C for AFC-77) and art as barriers which pinthe grain boundaries [8.20]. In AFC-77, the grain size can also be controlled throughaddition of refining elements such as niobium. With an Nb content of about 0.2%, grainsizes between 4 and 18 {im were obtained for material initially austenized at 1700;F

90

- . 1 M M i s s \ i . \ K - \ c i s r , s u n s I M I I S K U M I I M < I K < , I

(92> C | a n u subsequent ly re-austcni/c<i at t empera tu res be tween KidO and 2200 F (870and 1200 C l . If prior a u s i e n i / a t i o n is etlected at 2200 F (1200 Cl . a large •Train size willbe o b t a i n e d independently of the r e - a u s t c m / m g t e m p e r a t u r e i-S.V).

K.5. Ageing Reactions

A g e i n g r e a c t i o n s in t h e s t e e l s c o n s i d e r e d h e r e a r e e s s e n t i a l l y c o n t r o l l e d b \ ;|-,e s t r u c t u r e -of t h e l a t h m a r t e n s i l e o b t a i n e d a f t e r t h e s o l u t i o n t r e a t m e n t a n d iv. i h c n a i u r e a n dc o n t e n t o f t h e a l l o y i n g e l e m e n t s p r e s e n t . T h e d e f e c t - r i c h . m i e n ' s t r u c t u r e of i l , c m a n e n s i t eg o v e r n s n o t o n l y t h e k i n e t i c s o f t h e p r e c i p i t a t i o n r e a c t i o n s ( , ; . C h a p t e r 2 . S e c t i o n "• > ">. ib u t a l s o t h e m o r p h o l o g y o f t h e p r e c i p i t a t e s a n d e v e n t h e t v p c .,f p n a - c f o r m e d . , - b>f a v o u r i n g p r e c i p i t a t i o n of e q u i l i b r i u m r a t h e r t h a n m e t a - U i h l e c o m p o u n d - | h e m t e r -m e t a l l i c p h a s e s t h a t c a n f o r m in s t a i n l e s s m a r a g m g - t e e l - a n d a l l o \ - wi l l m m h e r e M e w e dw i t h t h e s i m p l e s y s t e m s s e r v i n g a s a n i n t r o d u c t i o n I > t h e m o r e c o m p l e x o n e s u p i L . t i ofh i g h - s t r e n g t h c o m p o s i t i o n s .

S.5.I. Inm-Chrnmium System

The present general consensus i- that the binary Fe-Cr system exhibits a miscihilny napat temperatures nelow «33O F (500 Cl. with the b.c.c. solid solution decomposing into nvoisomorphous phases, one rich in chromium u ) and the other rich in iron u i . Figure S.10shows the latest equilibrium diagram available [X.2I]. which incorporates resulis\if threebasic studies [S 22 to K.2-4] concerning the i-phase range. Although the data pertainingto the euiecloid decomposition temperature of this phase and to the boundaries of thevarious fields are still controversial, the diagram does indicate that steels with chromiumcontents of 10",, and more can undergo a decomposition reaction on ageiim at temp-eratures below 930 F (500 C). It has actually been shown [ti.25] that hardening of ihe17-4 PH steel (l6",,Cr) is due to the formation of a chromium-rich ferrite U i: the samephase lias been identified [A'.IVi] in a 12",,Cr steel subjected to fatigue stressing at 500 F(260 C). Similarly, several investigators [H.27. S.2K] have shown that ageing for relativelyshort periods at SSO F (475 Cl is sufficient to cause precipitation of the r phase in Fe-Cralloys containing 20",,Cr. Finally, the structure of an Fe-46Cr alloy aged at V30 I (500 Clfor 9650 hours has been found to be free from a and to consist solely of the y and v.phases [8.29].

On the other hand, the a-phase has been observed [8.17] in thin foils of Fe-Cr alloys afterageing at temperatures as low as 660 F(35OC). This would tend to show that decom-

Fig. 8.10.Iron-chromium phase diagram.After R.P. ELLIOTT [8.21].

100010

Cr CONTENT, wt .7 .

.30 40 50 60 70 SO 90

40030 40 BO SO 70

Cr CONTENT, at ".'.

91

i i > H \ I I i d S i ' l M M i 111< . H •>'. V.! N : • 1 11

el I U O ' F I M M O-Zhldark-fidd illumination I •/) 1110 F <600"O - 2000b 30,000

Fig. 8.11. — Thin-foil transmission electron micrographs of a13C'r-20Co steel after ageing. After D. COLTSOLRADIS el al. [8.31].

position of the Fe-Cr solid solution into :c and y. is not the final stage but an intermediateone, the equilibrium product, viz. the c phase, being seldom formed because of its muchmore complex structure.

8.5.2. Fe-Cr-Co and Fe-Cr-Ni Systems

The occurrence of a four-phase equilibrium reaction between s, y (f.c.c), a (b.c.c), and£ (h.c.p.) at 1110'F (600"C) in the Co-Cr-Fe system was reported in 1959 [8.30].Simultaneously, 78Fe-12Cr-IOCo was given as the composition of the triple point betweenthe a-ferrite field and the three-phase fields, v. -\- y • z and a -r e — c. at 1110"F.According to these data, steels containing about I2",,Cr and IO",,Co should not be stablebelow 1110 F. However, no equilibrium sections of the diagram below HIO'F werepresented in the above-mentioned work, in which no account was thus taken of theexistence of a decomposition reaction in the Fe-Cr system. In fact, identification of thephases forming in ternary Fe-Cr-Co alloys after prolonged ageing sets a difficult problem,which was tackled only recently.

After ageing a steel containing 13°,,Cr and !0%Co at 930'F (500'C), a finely dispersedprecipitate forms which is not a Cr-rich ferrite and which, even after 2000 hours, couldnot be identified as a a phase [fi.3i]. This precipitate is probably an intermediate phase,the term " intermediate " being intended to reflect the difficulty involved in identifyingaccurately a fine dispersed phase r?.ther than to designate a phase with a specific crystal

92

STAINI I.SS MARAOING SIKIiLS — PHYSICAL METALLURGY

structure. Similar behaviour is observed on ageing a l3Cr-20Co steel at 930 F (500 C)although, in this case, the " intermediate " phase occurs sooner and is denser than in the10",,Co steel (Fig. 8.11 a and h). The fine uniform distribution of the precipitate indicatesa homogeneous-type precipitation and reflects the effect of the high dislocation densityof the martensite.

On ageing at 1110 F (600 C), no precipitate is observed in the l3Cr-10Co alloy; in the13Cr-20Co alloy, owing to the increase in the degree of supersaturation of the solidsolution on account of the higher cobait content, formation of the T phas'j. observed ufter2000 hours, is preceded by homogeneous precipitation of the "" intermediate " phase(Fig. 8.11 c and </). This provides an indication that the same processes are involvedat 930 and 1! 10 F.

Similar metallographic studies on a marlensilic 13Cr-6.5Ni steel aged for 500 hours atI !10"F (600 C) failed to reveal the formation of any precipitate [H..V]. This was tenta-tively attributed to an effect of nickel on the fieid boundaries in the Fe-Cr system,particularly in depressing the solvus of the n phase or of the chromium-rich * phase.However, electrical resistivity measurements have provided evidence for the occurrenceof similar decomposition phenomena in Fe-Cr-Ni and Fe-Cr-Co martensites [S.I7].Figure 8.12 shows the temperature dependence of this property for an Fe-Cr-Co and anFe-Cr-Ni steel in two heat-treated conditions. For the heating rate considered (5 F/min.i.e. 3 C/min), decomposition of the solid solution in the Fe-Cr-Co alloy occurs attemperatures between 840 and 1200"F (450 and 650C); in the Fe-Cr-Ni martensite. u isrestricted to temperatures below 930 F (500=C). The decrease in resistivity occurringabove 1200 F in both steels is due to the transformation of ferrite to austenite: iiic extentof this reaction appears to be larger in the Fe-Cr-Co steel, probably due to its higherchromium content. Ar. regards the product of the decomposition, it should be noted thatsignificant amounts of a phase were identified only after prolonged ageing : 1000 hoursat 1110 and 1290 F (600 and 700*0 for the 19Cr-!5Co steel, and 1000 hours at 930 F(500 C) for the 14Cr-8Ni steel.

The hardening response of Fe-Cr-Co alloys, although appreciable, is not as high as thatfound in more complex alloys containing molybdenum. Typical age-hardening curves

TEMPERATURE, °C400 KM

— 1920°F(10501'C)-1h,A.C.

- o - DITTO + 930•F(500°C)-tOOh,AC.

TFMPERATURE,°C400 600

-1920°F(1050°C)-lh.A.C.

- DITTO+750°F(400"C)-100h,A.C.

E1.0C

osg

08 K

0.7 E

0 . 6 "

100 600 1200T E M P E R A T U R E , ° F

1600 400 BOO 1200T E M P E R A T U R E , " F

Fig. 8.12. — Temperature dependence of electrical resistivity of Fe-19Cr-15Co and Fe-14Cr-8Nisteels in the annealed and annealed + aged conditions. After K. BUNOARDT el at. [8.17].

93

I O H \ i l - l O \ l M M M i H I C i M - S I K l S C i I H S I M I S

HOLDING TIME AT INDICATED TEMPERATURES, hours

1 m M.l 5 1 tftxt i'l" . igeina t e m p e r a t u r e , ind tune on h a r d n e s s of r e - l l > l r - l ••(>' a n d I f - I 4 ( Y UN] slccls.Al ter K. B L Ni.AKiir t'i at. [.v/."]. I n t u a l ' - a r u l m o n : 1920 I ( 1050 (_) - 111. A C.

are reproduced in Figure X.I3 for the Fe-l9Cr-15Co alloy discussed above, as well asfor the Fe-!4Cr-!<Ni composition. The hardening response is quite rapid, with virtuallyno incubation, and the hardness maximum is shifted to lower times as the temperatureis increased: peak hardnesses are generally lower in the Fe-Cr-Ni than in the Fe-Cr-Costeel. The 0.02',, offset yield strength of Fe-Cr-Ni steels quenched to lath martensitewas shown to increase for tempering temperatures up to 840 F (450 (.'). at which pointaustenite reversion occurs f#./5]: the observed increase was related to the recovery ofstresses generated b\ the martensite transformation.

From the aboxe coi.siderations, it would appear that in Fe-Cr-Co, and possibly Fe-Cr-Ni,martensitic steels the precipitation reactions that occur during ageing are based primarilyon the decomposition of the Fe-Cr solid solution to n phase, although the latter has onlybeen positively identified in overaged samples.

8.5.3. Fe-Cr-Co-Mo ami More Complex Alloys

The simultaneous presence of cobalt and molybdenum in stainless steels has a pronounced" synergistie " effect as is evident from many investigations [8.6, 8.11, 8.16, 8.17, 8.32,8.33]. As an illustration, Figure 8.14 reproduces the change in hardness cf 13%Cr steelsaged for 128 hours at temperatures between 750 and 1110'F (400 and 600 C) as a function

200SO 80 BO 120 K0

WL'/. Co.wi.% Mo

Fig. 8.14.— Hardness of Fe-13Cr-Cc-Mo steels aged for128 h at indicated temperatures as a function of%Co x %Mo product. After E. DlDERRlCH el al. [8.34].

TABLE 8.4. — COMPARISON OF LATTICE PARAMETERS OF INTER-METALLIC PHASES LIKELY TO OCCUR IN Fe-Cr-Co-Mo ALLOYS

After W.B. PEARSON [8.37]

Phase

cr-FeCr(J.-Fe7Mo6

FejMo LavesR (Mo-Co-Cr)/ (Fe-Cr-Mo)

System

TetragonalRhomb.-hex.HexagonalRhomb.-hex.B.c.c. (a-Mn)

a, A

8.79954.7464.74

10.9038.92

c,A

4.544225.78

7.7319.342

cla

0.51645.4320.6131.774

94

X. S! AINil IK1-. MANAGING S I I I . 1 S IMIVSH Al M M A I I L K G Y

of the (« l . " ( . Co • \vt.",, Mo) product, ll is seen thai, ut all the ageing iemperaturesconsidered, the hardness increases with increasing (Co • Mo) product. The time toproduce peak hardness in such steels decreases as the ageing temperature is raised from750 to 1200 F (400 to 650 C) [8.34. 8.35].

A considerable number of inlermetallic compounds are likely to form during ageing inl f -Cr -Co-Mo steels depending on their composition and ageing conditions. That thisis so has been made clear in a recent review [SJfi] of the equilibrium phases present inthe quaiernav> diagram and its associated binary and ternary ssstems. Table K.4 ei'.esthe lattice parameters of some of these phases.

Table 8.5 summarizes the literature data concerning the phases thought to be responsiblefor hardening in cobalt-containing stainless maraging steels. As can be seen, a \arieiyof compounds have been quoted. ll should be borne in mind, ho\u"\cr. that because ofiheir structural similarity, their positive identification has. in general, onh been possiblein overaged specimens [<S.35, H.40\. In carbon-containing steels, carbides of the M->C.M;iC,, and M,,C types were identilied. while the ^-Ni-.Ti phase was obsened in titaiiiuni-modifled compositions.

Tiie resulis of investigations to determine the compositions of the extrar.ed phases areincluded in Table 8.5. The chromium content of the precipitates is much higher, and the

TABLE 8.5. — SUMMARY OF PHASES IDENTIFIED IN MARAGING STAINLESS STEELS

Ref.

18.31]

W.-il)

l».3l]

IS-'-'1[8.38]

[8.39]

[8.39]

[8.40]

W-13]

[i.35]

[8.33]

[8.41]

[8.6]I

1Composition

(\vt. "„, bal. Fe)

13Cr-10Co13Cr-20Co

13Cr-20Co-2Mo

I3Cr-20Co-5MoDitto

13Cr-13Co-4Mo-5Ni

14Cr-16Co-5 M o-0.5V-0.002C(low-carbon AFC-77J

14Cr-14L.o-5Mo-0.5V-0.17C(AFC-77)

12Cr-10Co-6Mo-0.1C

10Cr-I0Co-6Mo-7Ni

10Cr-13Co-4.5Mo

17Cr-8Co-3Mo-8Ni-0.02C

12Cr-14Co-5MooNi-0.5Ti-0.lAl12Cr-SCo-2Mo-SNi-0.8Ti-0.02C

Conditioni

930 F (500C) - 2000h930 F (500 C)-2000h

1110 F (600 C)-2000h930'F (500 C) - 2000h

1110 F (600 C)-500hIIIO'F (600O- 500h1110°F (600*C) - 2000h93O°F (500°C) - lOOh

1200°F (650°C)-2 h 2h1290DF (700°C)-2 + 2ht400°F (760°C) - 2 + 2h1500^ (815"C) - 2 4r 2h1200°F (650'C) - 2 + 2h1290°F (700°C) - 2 4- 2h1400°F (760cC) - 2 4- 2h15WTF (815=C)-2 4 2h1020°F (550' C) - lh1110°F (600X) • ihi200°F (650'C)- lh1290F (700X)- ih975°F (525°C) - 4h

1UO°F (600°C)-4h93O°F (500°C) - 1200h

11IO°F (600'C)- 1200h1290'F (700°C)- 1200hcold worked 97% +

84O"F (540°C)- 17h1560°F (850°C) - lh880-nF (475°C) - lOOOh

Phases

intermediate phaseintermediate phaseintermediate • aintermediate - ainiermediate • aa ^ FeiMoa 4- FeiMoclustering —> (Fe.CrhMo

Fe^Mo -i- y (traces)FeiMo + x (traces)X -r Fe2Mo (traces)X + FeiMo (traces)MzsCo + FciMo 4- x (traces)M23C6 4- Fe^Mo 4- /M6C + y + M,3C6M6C 4- x + M23C6

M-.CM2C r RM-.C -t- RM 6 C + R

RRRRX

a - x

a + x (grdin boundaries)7)-Ni3Ti + Fe2Mo 4- x

Ccof intCr

40.5

34 5

31.543.1

22.117.414.717.1

—

25.6

—

27.623.224.4

—

impositiceriVietalli

Co

15.S

16.8

17.31.9

(5.6NO10.3!0.19.09.9

1.8

—

6.08.0

10.0

:

—

>n (wt.°c precipi

Mo

_

I

3.3

6.9

28,0

29.232.034.625.5

27.8

31.627.58.1

:

—

,)tates

Fe

42.7

43.7

44.221.5

2S.32S.240.746.7

—

44.8

—

34.941.457.4

:

• —

95

I ; > R \ U - l ONT-UNINU HICH-STKhM. I H STKELS

.jl sl.lt> F I 'OO ( I - IJ f 'h • 80 .000 M 9 J 0 ° F • >0O C> - 2 0 0 0 h • 80 .000

Fig. 8.15. Thin-fni! transmission ciccirou micrograph ol a I3Cr-20Co-2Mo steel after agciny. Afier D. Coi TSOI RADIS el al. [SJI\.

cobalt content alwavs lower, than those of the parent steel. After ageing at temperaturesof the order of 1110 F (600 Ci, the precipitates are somewhat richer in molybdenumthan the base metal, although not sufficiently so to indicate the formation ol molybdenumimermeiallic compounds such as Fe^Mo or Fe7Mo(,. Ageing at 1200F (650'C) or abovegives rise to a very marked enrichment of the precipitates in molybdenum, whichsubstantiates the occurrence of the Fe^Mo. R or y phases.

Intermetalliv compounds. In Fe - 13Cr - 10 or 20 Co alloys the addition of 2 or 5",,molybdenum does not modify qualitatively the behaviour of the base composition asdescribed in the preceding section [8 31]. Ageing at 930 and 1I1O=F (500 and 600 C)gives rise to the precipitation of an " intermediate " phase and eventually to that of n(and of Fe^Mo in the case of the 5%Mo p'lo"). Figure 8.15 shows the morphology of theprecipitate formed in the 2",,Mo steel after ageing at 930'F (500;C). The precipitatereaction during ageing at this temperature is probably preceded by clustering of soluteatoms, mainly Cr and Mo, which is facilitated by the defect-rich microstructure. Theclustering mechanism is possibly similar to that found in Fe-Mo alloys [8.42]. Suchclustering, followed by the nucleation and growth of Fe2Mo-type particles, has beenobserved in an Fe-13Cr-13Co-4Mo alloy aged at 930F (500'C) [8.38]. Nucleation ofparticles on dislocations rather than on martensite lath boundaries seems to be a generalfeature of precipitation reactions in these systems.

In both AFC-77 and a low-carbon version of this steel, significant amounts of the FeiMoLaves phase were observed after ageing in the temperature range 1110-1290'F (60O-70Oc>C)[8.39). The presence of an Fe2('Mo,W) Laves phase was also established in various 12%Crsteels after long-time exposure at 1110 to 1470°F (600 to 800°C) [8.43].

Formation of the R phase was reported [8.40] to begin at 1110°F (600"C) in a 0.1C-12Cr-!0Co-6Mo steel, although its presence was only positively established by single-crystalelectron diffraction analysis of samples treated at I200°F (650sC). The composition ofthis phase (Table 8.5), which was still detected after short-time ageing at 129O°F (700°C),is very low in cobalt (1.8 %); this tends to show that it is not of the conventional Co-Cr-Motype, but rather of the Fe-Cr-Mo type. It should be noted, however, that the existenceof an R phase in the Fe-Cr-Mo ternary system has not been reported [8.44]. The R phasewas also observed [8.13, 8.35] in other Cr-Co-Mo steels after ageing at tnode>4lc!y hightemperatures (see Table 8.5).

For higher ageing temperatures or longer ageing times, the stable phase appears to be -/_,but it has not been possible to establish whether its occurrence results from the trans-

96

2000

AINIFSS V1ARAGING STEF.LS — PHYSICAL METALLURGY

TEMPERATURE, °C200 400

Q6 OB 10TITANiUM CONTENT, w.t.%

M 200 600 1000TEMPERATURE, °F

1400

l-'ig. 8.1ft. Effect of t i tanium content on yield and tensile strengthof Fe-12t"r-(tNi-5Co-2Mo-0.02C steel. After K. Bisc.ARDTctal. IX.I7).Condition: ISfiO F (S5Q C) - 30 mm. W.Q. • 900 F <4H0 C l - d h . A . C .

Fig. 8.17. — Hardening response of 14Cr-14Co stainless steels as affect-ed by Mo and C contents. After E.J. D u . i s and L. HABRAKIN [t'.S].

Initial condition : 2000 F (1095 C ( - In. O.Q refrigerai.on.

formation of a phase formed at lower temperature, e.g. R, c or Fe2Mo. For instance,the 7 phase was identified in a 10Cr-13Co-5Mo steel [8.35] after ageing at I290T (700 C)for 1200 hours. Similarly, the y phase was reported to replace progressively Fe;Mo inAFC-77 on ageing at 1290'F (700;C) or above [8.39]. whereas in Pyromet X-15, thestrengthening phase was stated to be R(Co-Cr-Mo) and,or / 18.45]. Cold reductionprior to ageing apparently extends the range of formation of this phase to iowertemperatures, at which it occurs in association with a [8.33].

The merit of using tungsten as an aiioying element in Fe-Cr-Co steels was evaluated byseveral investigators [8.11, 8.16]. The hardening effect of this element, though appreciable,is lower than that of equivalent at. % additions of molybdenum. Hence tungsten is notincluded in any of the commercial stainless maraging steels listed in Tables 8.1 and 8.2.

Titanium has occasionally been used as a hardening addition, particularly in steels(.-•'itaining nickel such as Ultrafort 401. Its effect in a 12Cr-8Ni-2Mo-5Co steel is illustratedin Figure 8.16. Ageing a steel of this composition for 1000 hours at 890:F (475 C) [8.6]results in the precipitation of Tj-NijTi, Fe2Mo and small amounts of / at dislocations.

Copper additions (2 and 5%) to a 14Cr-20Co alloy cause appreciable hardening [8.12].However, transmission electron microscope examination of copper-containing Cr-Co-Mosteels aged at 930 and 1110°F (500 and 600°C) failed to reveal any copper-base phases.

Carbide reactions. Some cobalt-containing high-strength stainless steels, e.g. AFC-77and Pyromet X-12, contain carbon in non-negligible amounts. The data available onthese steels show that precipitation of carbides and of intermet'allic compounds overlap,and that the hardening effects of both species are additive and complementary.The hardening responses of three chromium steels, one virtually carbon-free and theother two with substantial carbon contents, are compared in Figure 8.17. The carbon-freeCr-Co-Mo steel starts to harden at about 660°F (350°C), and reaches peak hardness at1020°F (550°C). The molybdenum-free Cr-Co-C steel shows a secondary hardness peakat about 885°F (475CC), which is followed by a rapid decrease in hardness. The steelcontaining both carbon and molybdenum exhibits a behaviour which is the summationof the two preceding ones. Thus, it is seen that strengthening of Cr-Co-Mo-C steels ismainly by precipitation of carbides for ageing temperatures up to about 840°F (450"C),and of intermetallic compounds in the 930-11IO°F (500-600"C) range (though carbidesmay still play a minor hardening role for tempering temperatures above 840°F) [8.39].

97

I'OK.-UT-'.ONI-UNINU HK.HSIRENGTH STEELS

TEMPERING TEMPERATURE.'C

so. ap *-^-*^.6oo- I

•4" ^ * > ^ ^

f i e . S.iH. ThiP.-Ioii t r ansn : i*s ion e lec t ron m i c r o g r a p hof M L - after d o u b l e t e m p e r i n g I - 2 hour s ) at IOL*> (-

(?wo Ci . \fter 1. D I I I I K H K M <•/ ill. It.II}. I7.51K)

TEMPERING TEMPERATURE, "F

I-ig. S.I»*. ! Meet of t e m p e r i n g t empe ran i r e on R . T . f rac ture (nughm1-.*, 0 .2" , ,V*S arul n o t c h e d tens i le s t reng th ol" A K - 7 7 . After R . T . A I I . T it til. \S.7\.Initial c o n d i t i o n : 14(11)1- | 7 M > C i - 2 h • 1 W 0 F I !(M0 Cl - Ih." ( ) (J.

• r e f r i ge r a t i on . 2 2h t e m p e r s .

The carbide reactions that occur in AFC-77 during ageing have been identified [li.39]as follows (see Table S.5) : M;-.Cn carbides precipitate on tempering between 1200 and1500 F (650 and 815 C). v.liile MftC cat bides start to occur and finally become predominantin ihe 1400 to 1500 F 1760 to 815 C) temperature range. Figure 8.18 reproduces themicrostructure obsened in AFC-77 (the Cr-Co-Mo-C steel in Figure 8.17i after temperingat a temperature just above that giving rise to peak hardness.

In the steel containing 0.1C-12Cr-10Co-6Mo (see Table 8.5). M;C carbides precipitateas needles at from 1020 to 1200 F (550 to 650 C) [8.40\. There is evidence thai suchcarbides form in steels of this type at even lower temperatures, e.g. 930 F (500 C) [8.16].In the ••••'•! just referred to. M;C gives way to MhC on tempering at 1290 F (700 C).

In summary, the carbide reactions in higher-carbon stainless maraging steels seem toproceed as follows : M^C is preponderant for tempering temperatures up to- M 10 F(600 Ci. and can thus be assumed to be mainly responsible for the hardness peak associatedwith secondary carbide precipitation (Fig. 8.17): M^Q, carhides form at 1110 F (600 Cland above: beyond about 1290 F (700 C). MhC carbides become progressively pre-dominant until soluiioning of all the precipitated phases occurs. This p ittern is similarto that prevailing in 12"oCr steels [8.1].

8.5.4. Concluding Remarks

Stainless maraging steels are normally aged in the 840-1200 F (450-650C) temperaturerange, depending on the properties required and the intended use (see Chapter 9,Section 9.1). The nature of the matrix is an essential feature in controlling the ageingprocesses occurring in these steels. In this respect, cobalt plays an important role sincethis element is capable of producing a martensitic structure in steels with the highchromium contents needed for corrosion resistance. The defect-rich martensite accountsfor the lack of incubation and the high reaction rates encountered during ageing.

At low ageing temperatures (750 to 930=F, i.e., 400 to 500vC), a finely dispersed phase isobserved in Fe-Cr-Co alloys, and is presumably related to the decomposition of the Fe-Crsolid solution. This reaction progresses beyond the stage of formation of the chromium-rich y. phase towards formation of the o- phase, although the latter has not been identifiedwith certainty in such alloys, because of its fineness. At higher ageing temperatures, theprecipitate coalesces, permitting positive identification of the er phase. The early stagesof precipitation may be preceded by clustering followed by formation of zones.

98

X S1AINU SS VARAGINC, STFLLS PHYSICAL METALLURO

Molybdenum, when present, may also participate in the clustering and zone-formationstages of precipitation. For low,molybdenum contents, up to 3.',,, lhe phases formed areof the finely-dispersed i-lype [8.3!. 8.33]. For higher molybdenum contents. \i:.. from4 to 6",,, this element leads to the formation of other phases : Fe2Mo. R or / . [he latterprevailing at the highest ageing temperatures. The composition of the precipitates tendsto confirm the occurrence of Fe^Mo and / phases, or 'H- formation, under specilicconditions, of an Fe-Cr-Mo R phase.

As regards the enhancing effect of cobalt on the precipitation reactions, several inter-pretations have been proposed. Cobalt may decrease the solubility of chromium [8.32]in addition to lowering that of molybdenum (see Chapter 5, Section 5.3.3). It mav alsoraise the r.acleation rate h\ increasing the degree of supersaturation of the solid solution[8.32]. by stabilizing the nucleating dislocations [8.40] (see also Chapter 3. Section 3.<->.6)or. as stated earlier (Chapter 5. Section 5.3.3). by decreasing the stacking-fault energv ofthe matrix, thus increasing the number of dislocations available as sites for precipitation.Finally, cobalt may induce ordering of the matrix through an enrichment effect resultingfrom its low solubility in the precipitates [8.32, 8.40].

8.6. Strength ' Toughness vs. Structure Relationship

A considerable amount of work has been carried out with the aim of improving thetoughness of AFC-77: although some of the numerical data will be used in the nextchapter, it is appropriate to summarize here the more fundamental aspects of the problem.iL'. the relationship between toughness and tensile properties, on the one hand, andstructure, on the other.

8.6.1. General

The effect of tempering temperature on the yield strength, notched tensile strength andplane-strain fracture toughness of AFC-77 is shown in Figuie 8.19. The notched strengthand toughness are excellent up to 700 F (370 C). but both these properties exhibit a sharpdrop between 700 and 800 F (370 -and 425C). Other data [8.8] situate the onset of thedecrease in fracture toughness at about 800:F. On the other hand, measurements of themergy absorbed in impact tests on notched specimens show that the rapid loss inioughness takes place on tempering between 600 and 700cF (315 and 370 C) [8.7].

Fractographic analysis of centre-notched tensile specimens revealed [8.7] t!;at the brittlebehaviour is characterized by cleavage fracture and not by a predominant intergranularor grain-boundary fracture. Examination of replicas failed to allow identification of theprecipitation reactions occurring in the 600-1000°F (3I5-540DC) range; only after temperingat 1100 and 1200°F (590 and 650=C) were M23C6, M6C and Fe2Mo precipitates identified.On the basis of earlier work [8.39], which is in agreement with the preceding observations,it was concluded [8.7] that Fe2Mo is a more likely cause of embrittlement than are M^Când M&C (in Figure 8.17, the onset of hardening in the carbon-free steel lies at about660°F, i.e., 35O°C). It is the present authors' opinion that, considering the temperatur:level at which embrittlement is observed, the process involved is probably either thedecomposition of the solid solution or a carbide precipitation reaction.

8.6.2. Effect of Retained Austenite

The effect of retained austenite on the fracture toughness of AFC-77 is illustrated inFigure 8.20, in which Kic has been plotted as a function of tempering temperature, withthe retained austenite content as an independent parameter. The beneficial effect ofretained austenite is evident. In this respect, it will be recalled that an optimum stableaustenite content of 5 to 15% was one of the basic considerations in the development ofPyromet X-23 (see Section 8.3.3).

99

\t T (.UNTUNING HIGH-STRlNCiTH STEELS

(Mi, MN/n'1S00 IOT1 2O30

50C• " 600 ' 0 0 600 900 1000 H0OG TEMPERATURE T

Pig. S ^0. Mltvl i'l relumed aiMi'nife on Iractureloughnev. of A K ' - " After I). U I K S I I R |>./wj

Condition : auê t r ' e J ai ! "00 to 2(00 h ( C 5M 1150 O i . r fig >»!>>. ; 2h temper

I i£. s.;i StroiBih toughncsi rela-Iior^hip lor -MC-""* and steels vienv-ed therefrom. After D. U I B S I I B |A'.v|.

The toughness increase produced by the ausienite is apparently due to the latter's crack-arresting ability. When a crack propagating through a martensitic region encounters anaustenite area, it may either continue into the austenite or trace a /ig-zag path throughinterlocking martensiie laths; both processes lead to an increase in the energy for fracture[#./9]. High-magnification cine-films of cracks growing in AFC-77 sheet ha\e shownthat these are in fact arrested on reaching an austenite area, arourd which they branch andpropagate as the load is increased [tV.yj.

In AFC-77. this mechanism is most efTecti\e in the 240,000 to 260.000 psi (1600 to1800 MN m-) strength range produced by tempering at 500-700 F Cf, ' 370 C); at thehigher tempering temperatures needed to obtain strengths in the 26O.OOU 'o 290,000 psi(1800 to 2000 MN m:) range, the amo int of austenite that can be retained decreasesmarkedly. So as to overcome this diffic ilty. the composition of the steel was modified byadding abnir. 1 "„ of the austenite-stabilizing element, nickel. After tempering at 800 F(425 C). the modified steel exhibited an unusually large elongation of 28",, for the260.000 psi strength level achieved. This has been attributed not only to the stability ofthe retained austenite, but also to the occurrence of a stress-induced martensite trans-formation [S.9]. AFC-260 is another nickel-modified version of AFC-77; however, in thiscase, the simultaneous decrease in the carbon content from 0.15 to 0.08 "„ tends to negatethe effect r.f the 2.0 %Ni addition [8.9].

The toughness of AFC-77, AFC-260, and " Alloy B " (the 1 %Ni grade just referred to,which also includes 0.2"oNb for grain refinement, see next section) is plotted as a functionof ultimate tensile strength in Figure 8.21. The steepness of the AFC-77 curve is a reflectionof the diminishing austenite content at the higher strength levels. In AFC-260, austeniteis present in large amounts up to the maximum strength of 265,000 psi (1850 MN/m2),so that the strength-toughness line is considerably less steep. " Alloy B ", which retainsaustenite up to an ultimate strength level of 290,000 psi (2000 MN/m2), has a strength-toughness line which is a continuation of that for AFC-260.

8.6.3. Effect of Delta-Ferrite

Delta-ferrite formation in AFC-77 and AFC-260 during high-temperature austenizingtreatment, which is desirable to produce significant amounts of retained austenite, reducesthe strength and toughness, especially after tempering at temperatures above 700°F(370°C) [8.10]. As stated in Section 8.2, removing this S-ferrite by re-austenizing at a

100

9. STAINLESS MARAGING STEELS — PROCESSING AND PROPERTIES

A.OE'NC- TEMB 450

y-r4GFING TEMPERATURE.'C

«X) «50 500 _S50_

soo 900 nooAGEIMG TEMPERATURE, 'F

SQO 900AGEING TEMPERATURE, "F

TENSII F STRFHGTH, MM/m21600 T700 1800 1900

1 tr. 1 V«700"F(370-C)\

VI20 240

TENSILE

•F(260-C); ;

V—î^-jooo'Rsa)

2605TREN3TH, 103 psi

1

-

_

C)_

20

W

8 0 |

6 0 *

40

2S0

Fig. 8.22. — Strength and toughness of low-carbon AFC-77as a function of auslenizing treatment and temperingtemperature (2 ~ 2h treatment). After D. WEBSTER [8.10].

Toughnesses were measured on sub-size specimensand are thus of relative significance only.

Fig. 8.23. — Effect of grain size on strength: tough-ness relationship for AFC-77. After D. WEBSTER[8.46]. Fine : soft-tempered - cold rolled;coarse : hot rolled; both followed by 1800'F(980°C) anneal, refrigeration. 2 -- 2h temper.

lower temperature results in a better compromise between strength and toughness atall tempering temperatures. This is illustrated in Figure 8.22 for sheet of a low-carbonmodification of AFC-77 (0.1I5%C); it is seen that the strength level after the two-stagetreatment is identical to that reached after the single-stage treatment at 1900:F (1040"C),but that the toughness is much higher. The superior toughness must be attributed to thedecreased S-ferrite content associated with a sufficiently high retained austenite content.

8.6.4. Effect of Prior Austenite Grain Size

It is known that prior austenite grain size affects the yield strength of stainless martensiticsteels through a Petch-type relationship [8.151 The effect of this parameter on the strength/toughness relationship was studied in the case of AFC-77 [8.46]. The method used toobtain grain refinement consisted in soft tempering the material at 1400 F (760C),followed by cold rolling to 50% reduction, and annealing at temperatures between 1800and 1900T (980 and 1040=C). This produced grain sizes comprised between 2.3 and60 (Am; the latter dimension is similar to that produced by standard austenizing withinthe above-mentioned temperature range. Refining the austenite grain from 60 to 2.3 u.mresulted in an appreciable strength increase at no cost to toughness (Fig. 8.23).

Another means of improving the strength/toughness combination in AFC-77 is to takeadvantage of the known grain-refining effect of niobium. A 0.2 %Nb addition proved tobe an efficient grain refiner; however, grain refinement was practically inoperative whenthe alloy was treated to ultra-high strength levels (<•/. Fig. 8.21, " Alloy A ") . As statedearlier, combining this addition with 1 %Ni (" Alloy B ") yielded excellent results.

9. STAINLESS MARAGING STEELS — PROCESSING AND PROPERTIES

9.1. Primary Processing

In view of their composition, air melting techniques can be used to prepare most stainlessmaraging steels. However, steels which contain fairly large amounts of reactive elementssuch as titanium (e.g. the Ultrafort grades) require vacuum melting. As is the case forother high-strength steels, consumable-electrode remelting may be of value, but the effectsof this melting process are poorly documented. Nevertheless, it is widely held that, forbest toughness and ductility in stainless maraging steels (particularly in large sections),

101

lOIUl l -u lMMMM. HIGH-STRENGTH STEELS

l \n\ i HOI WORKING OI SIA'NLLSS MARAGING STEELS

\K.'-'' [«./! Heal uniformly at 2050 F<l I20C >. Hot working temperature range is 1700lo 2050 F (425 to 1120 O . No particular precautions arc necessary oncooling. After cooling, anneai at MOO F (7h0 O to." 2 2 hours.

Pyromet X-12 («..] Heal uniformly to 2050 F (: 120 O. Immediately after hot working, largeforgings and billets should he sen! hot to a furnace and normalized at1 htxvi "Mi F (S7U-y;5 I I. air cooied to room temperature and then annealedat 1300 F (705"O for 1 l.i J hours and air cooled.

P>romct\-15 [0..?J Homogenize at 22OO-2W F ', I2O5-12M) C) for about 30min,in. of cross-section. Hot working temperature range is 1M10 to 2300 F (870 to !2M) O .Final working should be done from a furnace temperature of 2000 F (UW5 C).

Pyromet X-23 [v.4\ Forging temperature range is 1W0 to 2300 F (870 to 12h0 O . Final workingsho'.fd be done from a furnace temperature of 1900-2000 F 11040-1095 C).

the lowest levels of " residual " elements should be obtained. In pratice. this necessitatesat bast single \acuum melting.

The hot-working conditions recommended for several steels of interest here are shownin Table 9.1. Generally speaking, stainless maraging steels are ductile at normal hot-working temperatures, and their forgeability is similar to that of other high-strengthstainless steels.

The heat treatment of stainless maraging steels consists essentially of three steps, viz.solution annealing, quenching and ageing: for some steels {e.g. AFC-77) a refrigerationtreatment is recommended between quenching and ageing. The low-carbon grades do notnecessarily require a rapid quench for response to precipitation hardening. Typical hea:treatment schedules for representative steels are listed in Table 9.2. Depending on the

TABLE 9.2. — HEAT TREATMENT OF STAINLESS MARAGING STEELS

Steel

AFC-77 (annealed)

AFC-77 (aged)

»

AFC-260 (aged)

»

Pyromet X-12 (annealed)

Pyromet X-12 (aged)

Pyromet X-l 5 (afe?rt)

Pyromet X-23 (aged)

Pyromet X-23 (cast)

Llltrafort 401 (aged)

I.'itrafort403 (aged)

Heat treatment

1400 F ( 7 6 0 X 1 - : - : or 2 - 2 - 2 h. A.C.

1900 F (1040C) - lh. O.Q. / 500°F (260=C) - 2+2h, A.C.1900-F (1040 C) - In, O.Q. / —100F (—73 C) - lh / 700°F

<37OC)-2~2h,A.C.1900=F(I040'O- lh.O.Q./ 1100=F (595°C)- 2 + 2h, A.C.

1900 to 2000'F (1040 to 1095"O, In,O.Q./ — I00=F (—73cC)-8h / 700 to 800°F (370 to 425"C) - 2+2h, A.C.

1900 to 2000"F (1040 to 10955C)- lh,O.Q./--I00°F (—73"O-8h / 1000 to 1100°F (540 to 595°C) - 2+2h, A.C.

1300°F (705"C) - 2 to 4h, A.C.

1700'F (925'C) - Ih, A.C. / 9003F (480'C) - 2 to 4h, A.C.1700 F (925 C) - lh, A.C. / 1100F (595 C) - 2 to 4h, A.C.

1700°F (925°C) - 30min,O.Q. or W.Q.; 1025"F (550X)-4h,A.C.

1500 to I800'JF(81S to 980"C)-lh, W.O.O.Q. or A.C. / 900 to1050 F (480 to 565C) - 2 to 6h, A.C.

2100F(lI50'C)-4h. A.C. / 1700F (925°C) -4h. A.C. / 1O25=F(55OC)-4h,A.C.

1560 F (85OC) - 30 min, W.Q. / 880"F (475=CJ - 6h, A.C.

1560cF (85O°C) - lh, W.Q. / 895"F (480°C) - 6h, A.C.

Observations

For best machinability (Rc 37-40)

For maximum cold formabilityFor service temperatures below

700°F (370°C)For service :emperatures up to

1100T (595°C)

For high fracture toughness andgood stress-corrosion resistance

For maximum room- and high-temperature strength

For best machinability (Rc 30)

For maximum strength (Rc 50)For maximum toughness (Rc 38)

Rc50

R r 45-50 (Rc 29 after sol jtioning)

—

—

—

Ref.

{9.11

[9.1]

»

[9.7]

»

19.2]

[9.2]»

19..']

[9.4]

[9.4]

[9.5]

[9.6]

102

9 STAINLESS MARAGING STEELS - PROCESSING AND PROPERTIES

TABLE O.3. — TYPICAL ROOM-TEMPERATURE MECHANICAL PROPERTIES OF STAiNLESs MAKAUING STEELS

Steel •

AFC-77(aged at 700 Ft

AFC-77(aged at 1100 F)

AFC-260(aged at 800=F)

AFC-260(aged at 1000 F)

Pyromet X-12(aged at 900 F)

Pyromet X-12(aged at HOOF)

Pyromet X-15

Pyromet X-23

Ultrafort 401

Ultrafort 403

U

10' psi

255

290

224

254

236

182

235

258

242

245

T.S.,

MNII!-1

1758

2000

1544

1751

1627

1255

1620

1779

1669

1689

02°,

10' psi

200

214

188

228

210

150

215

237

227

242

Y.S,

A TV

„•:1.179

1476

1296

1572

1448

1034

1482

1634

1565

1669

Elong.,

"i;

17

10

20

14

13

15

17

15

!3

10

R.inA.,

52

32

56

44

38

50

60

58

60

60

C.V.N.impact.

ft.lb

18

9

—

—

—

11

15

18

41"

25»

J

24

12

_

IS

20

24

56'-'

34»

A'

10-'psi \M

62

23'

92

61

_

—

70

94

49-61

i

in.

68

25''

101

67

—

—

—

77

103

54-67

N.T.S.

U.T.S.

1.5(AV 3.9)

1.0(A', - 3.9)

—

—

—

—

1.4(K,= 10)

1.4(A' (= 10)

1.6(A', = 3.1)

1.5

Hard-ness,

Re

50

53

48

41

48

50

—

Modulusof elasticity.

10" psi

30

31

29

o,Vni1

207

214

200

Ref.

[9.1.9.7)

[9.1.9.7]

[9.7]

[9.7]

[9.2]

[9.3]

[9.4]

[9.6]

[9.6]

a Heat treatments as in Table 9.2; when several treatments are possible (AFC-77. AFC-260 and Pyromet X-12). they are identified by the ageing temperature.b According to DIN 50 115; DVM specimen (notch 3 mm deep. 2 mm radius).1 These figures are for steel aged at !000"K (54O°C).

application envisaged, a wide combination of properties can be obtained by appropriatemodifications of the heat treatment parameters. For example, ageing AFC-77 and AFC-260at 700-800 F (370-425C) optimizes fracture toughness and stress-corrosion resistance,whereas ageing at 1000-1100T (540-595 C) produces maximum room- and high-temperaturestrength. A special feature of Pyroinet X-23 is that, in the fully heat-treated condition, itexhibits a fair amount of retained or reverted austenite distributed uniformly throughoutthe matrix. The same structural characteristic can be obtained in AFC-77 and the AFC-260and " Alloy B " grades derived from it; it was shown in Chapter 8 (see Sections 8.2 and8.3.3) that the use of higher austenizing temperatures or of a double austenizing treatmentleads to an optimum austenite content and a minimum amount of S-ferrite in these steels.It will be seen in the next section that such treatments result in optimum con jinations ofstrength and toughness.

9.2. Properties

9.2.1. Strength/Toughness Relationship

Typical room-temperature properties are shown in Table 9.3 for several cobalt-containingstainless maraging steels. Yield strength levels generally exceed 200.000 psi (1400 MN/m-),except when a trade-off between this property and fracture toughness is sought (see, forinstance, the data for AFC-260).

The few fracture toughness data listed in Table 9.3 indicate that Klc values of 60,000 to70,000 psiVin. (66 to 77 MNm--^) , associated with ultimate tensile strengths of 255,000-260,000 pfi (1750-1800 MN/m*), are found in Pyromet X-23, AFC-260 and AFC-77. In

103

(. OH U i i O M -UMNli HKili SlKLNUllI S i l l IS

ULTIMATE TENSILE STRENGTH MN/mJ

_ ; "ALLOY A" j

^XPYROMET IX-23

Tig. **.I, Strength loughf l ivrelation>hip for \ a n o u s hiyh-

strcngth steeU.

\ H - " ausicni /vd .it 2IH10 i i l lAfter P . WlHSriK |V,V. VV|;

due It' Npeeimen configuration, values. i K n c — HXUK)0 psi \ in. ! cn j M beopuniiNlic.

.-. .: After D. W t n s r i H |<J./"OI( • - Allo> A "'I.

A F C - 7 " ausieni /ed at 1900 F i l l )n After E J . n- .Lis ei til. [9~lr. After D. WEBSTER [<JM\.

AFCOftO (III)After D. W H I S T I R

175 200 225" 250" 275 300ULTIMATE TENSILE rÊNGTH, 103 p s ,

'Alloj ! ) " ( • ) (IV)After D. VVfBSTiK 1«./OJ.

Ultrafort 401 and L'llrafon 403 afterH. BRANUIS and A. VON DIN STEIN>.S1 .6); Pyromct X-23 after Ref. [1.41,other steels after D. WEBSTIR IV.IO],

the former two. this favourable combination is attributable lo the appreciable amountsof austenite retained at the higher ageing temperatures used (975 to 1000 F, i.e., 525 to54XTC). Ultrafort 401, and AFC-260 aged at a lower temperature, exhibit A'k. valuesas high as 90.000 to 100,000 psi\ in. (100 to 110 MNm -3.2) at interesting strength levels.Strength and toughness of AFC-77 appear to be fairly isotropic [9.8]: the absence of adrastic drop in toughness in the transverse direction is indicative of uniform austenitedistribution.

The fracture toughness of various high-strength stainless steels has been plotted inFigure 9.1 as a function of ultimate tensile strength. In addition to the values givenin Table 9.3, this figure includes further data on AFC-77 and modified versions thereof,as well as on several cobalt-free high-strength stainless steels. Curve I pertains to AFC-77austenized at 2000T (1095 C) or higher [9.8. 9.9] and aged between 800 and HOOF(425 and 595'C); two points corresponding to the Nb-modified version of this steel (Alloy A)[9.10] also fall on this curve. Generally speaking, the toughness decreases and thestrength increases as the ageing temperature is raised. The shaded area designated IIapplies to the unmodified steel heat treated using the conventional 1900°F (1040°C)austenizing treatment [9.7, 9.9]. The marked beneficial effect of raising the austenizingtemperature is due to the presence of austenite, as discusst-d in Chapter 8 (cf. Section 8.6.2).Area III corresponds to AFC-260 austenized at 1900 F (1040 C) and higher, and agedat 800 to 1050cF (425 to 565°C) [9.9]; the data given in Table 9.3 [9.7] for this steel aremuch more conservative, and fall within the shaded area for AFC-77 austenized at I900°F.On the other hand, values even higher than the former have also been reported [9.10];

104

9 STAINLESS MARAGING STEELS — PROCESSING AND PROPERTIES

i2Dr—T2 :

AGEING TEMPERATURE,X400 500 C30 /00

AGEING TEMPERATURE, CC150 500 550"T

AGEING TEMPERATURE, °C500 525 550 5/5 600

a 2 8 0

" 8 0 0 BOO 12O0AGEING TEMPERATURE, °F

1«B 800 900 1000AGEING TEMPERATURE, °

950 1000 1050 1100 1150AGEING TEMPERATURE, "F

Fig. 9.2. — Effect of ageing temperature on room-temperature mechanical properties of several high-strength stainless steels.AFC-77: after F..J. DULIS el al. {9.7]. Initial condition: 2000 F(1095 C)- lh.O.Q. / —100 F (— 73 C) - 30min. Ageing time: 2 -2 hours.

Dotted yield strength curve is for specimens austenized at 1900 F (1040 C).AFC-260: after E.J. DULIS el al. [9.7]. Initial condition: 1900 F (1040 C) - 1 h.O.Q.; -100 F (—73 C) - 8 h. Ageing time: I 7 hours.Pyromet X-15: after [9J], Initial condition: 1700 F (925 C) - 30min. O.Q. Ageing time: 4 hours.Pyromel X-23: after [9.4]. Initial condition: 1700 F (925 C) - Ih.W.Q. Ageing time: 4 hours.

they form the first part of Curve IV, the continuation of which corresponds to Nb-and Ni-modified AFC-77 (Alloy B) [9.10]. Finally, the single data points available forUltrafort401 and Pyromet X-23 are favourably located with respect to the curves justdiscussed. On the contrary, the fracture toughness and ultimate tensile strength of PyrometX-15 and the cobalt-free high-strength stainless steels do not exceed 75,000 psi \ in.(83 MNm--'/2) and 245,000 psi (1685 MN/m2). respectively.

The effect of ageing temperature on the room-temperature mechanical properties of AFC-77,AFC-260, Pyromet X-15 and Pyromet X-23 is shown in Figure 9.2. The yield strengthtrough observed in AFC-77 aged at 1000cF (540°C) has been attributed to the fact thatthe ageing reactions which occur at this temperature cause the martensite to contract,thereby setting up internal tensile stresses in the retained austenite regions (whereas, atother ageing temperatures, these internal stresses are compressive) [9.70]. This troughonly occurs in steels that contain significant amounts of retained austenite, i.e., thoseaustenized at 2000°F (1O95°C); it is absent in steels austenized at lower temperatures.

The effect of cold working on mechanical properties has been determined in the case ofUltrafort 401. It has been found that this steel's behaviour is similar following cold defor-mation by drawing [9.77] or rolling [9.5]. Figure 9.3 shows the effect of cold drawing :there is a slight strength increase up to reductions in cross-section of 60%, but a very sharpone beyond this value (possibly due partly to the development of a fibre texture, thoughthis has not been ascertained): ductility decreases with increasing amounts of reductionbut does not drop to prohibitively low levels; retained austenite vanishes after 20% reduc-tion. Ageing of the cold-worked structure at 900°F (480°C) produces a strength increaseof the same magnitude as that observed after annealing; material deformed 98% bydrawing and then aged exhibits a strength of 355,000 psi (2450 MN/m2).

105

i OH \l I -i U N ! \ I M M - Mil .11 M Kl M . I II SI I 1 I S

50DEFORMATION,".

40 60DEFORMATION. %

I i\:. ' O I fleet "f cold work on tensile properties of L'lirafort 401. Afterk BiM.vkliT jnd W- Sr>Ri |v.^); Condition: 1 5M> h 1X50 Ci - 30 mm. W.Q.; colddeformed b> drjv.mil tdolted lines); aged j ! 900 h (4X0 f i for 6 hours (lull lines!.

Ttv: posMhihtv of improving the properties of stainless nvraging steels through thtrmo-tnechamcal treatment has also been investigated. The effect of strain ageing on theproperties of Lltrafort 401 was dealt with in the preceding paragraph. Figure 9.4 showsthe relationship between strength and toughness for conventionally treated, strain-agedand ausformed AFC-77 steel plate [V.o']. Strain ageing involved !0n

(, cold work after aninitial 500 F (260 C) temper: ausforming involved 50°,, deformation by rolling at 1000 F(540 Ci in a series of 10 passes. Strain ageing results in improved strength toughnesscombinations for final ageing treatments up to 800 F (425 C); ausforming produces amore modest improvement for final ageing temperatures of 500 and 700 F (260 and 370 C).The effect of much higher percentages of cold deformation (by drawing), combined withintermittent ageing treatment, on the properties of AFC-77 wire is shown in Figure 9.5.Ultimate tensile strengths as high as 600,000 psi (4150 MN/m:) were achieved on wirecold drawn from 0.05 in. (1.27 mmi initial to 0.0046 in. (0.12 mm) final diameter (92",vreduction), while values of around 500.000 psi (3450 MN.'m2) were still attained in wirewith diameters as large as 0.05 in. (1.27 mm), drawn from rod 0.25 in. (6.3 mm) in diameterand aged appropriately [V.I 2. 9.13].

200

, ion

GOOVELD STRENGTH. M

1600 2003

500'FCSO'C) • 9OO'FK6O-C)7OO'F07a'C) a IOM"F6WC)aOO'FiaS'C) a nOO"F<595-C)

V » V . STRAIN A 6 E D -

CONVENTI0NAL

200

150 !

100

?00 2S0VELD STRENGTH, «P psi

- 1500

40 60 80 TOREDUCTION IN DIAMETER, V.

Fig. V.4. — Strength/toughness relationship for AFC-77in the conventionally heat-treated, ausformed andstrain-aged conditions. After D. WFBSTER [9.8].

All specimens austenized for Ih at 2000°F (1O95X), oilquenched, refrigerated at —100'F (—73 C), and tempered for2 - 2h at indicated temperature. Ausforming (50% reduction byrolling) was carried out at 1000'F (540°C) prior to oil quenching.Prior to strain ageing (10% reduction by cold rolling), therefrigerated material was tempered for 2h at 50CF (260°C).

Fig. 9.5. — LTect of cold reductionby drawing combined with intermittentageing at 900'F (480 C) on tensilestrength of AFC-77 wire. After V.K.CHANDHOK and A. KASAK [9.12\.

Initial condition : wire 0.05in. (1.27 mm) in diameter aus-tenized at 2000°F (1095°C)and tempered at 500°F (260°C).

106

STAINLESS MARAGING STEALS - PROCESSING AND PROPERTIES

100

TEST TEMPERATURE, "C

.200 300 400

100 500 700TEST TEMPERATURE, °F

1B0

i Ultimate tensile strength vs. test temperature.Heat treatment as in Table 9.2 (AFC-77 andAFC-260 aged at 1100 and 1000DF. respectively).

O AFC-77After E J . Duus andL. HASRAKEN [9.75]

AFC-260Mod. 12%CrPH steelsPyrometX-15. After [9.3]Pyromet X-23.Ultrafort 401Ultraforl 403

After [9.4]After H. BRANDIS and

A. VON DEN STEINEN [9.6].

TEST TEMPERATURE. °C_ 0 150 300 iSD 600 -J5Q 0 I5C 300 450 600

-200 200 60) tOO -200 0 200TEST TEMPERATURE, *F

600 1000

b) Tensile and impact properties of Ultrafort 401 andUitrafort 403. After K. BUNGARDT md W. SPYRA [9.5] andH. BRANDIS and A. VON DEN STEINEN \9.6\, respectively.

B0TEST T E M P E R A T U R E ^

200 a n 400 500 600

,-IOOOS

200 W 600 800 BOOTEST TEMPERATURE, °F

1200 K00

c) Ultimate tensile strength of AFC-77 wiredrawn and aged to 550.000 psi 13800 MN/m :).After E.J. Duus and L. HABRAKEN [9.13].

Fig. 9.6. — Effect of temperature on the mechanical properties of several high-strength stainless steels.

9.2.2. High- and Low-Temperature Properties and Thermal Stability

The effect of temperature on the ultimate tensile strength of several steels considered inthis chapter is illustrated in Figure 9.6 a. Their temperature capability is generally superiorto that of the best cobalt-free precipitation-hardening steels. The temperature dependenceof the mechanical properties of Ultrafort 401 and Ultrafort 403 are shown in greater detailin Figure 9.6 h. Both steels are seen to become brittle at — 321 F (—196 C); the rapiddecrease in strength which they exhibit in the 840-1110 F (450-600C) temperature range,is due to austenite reversion [9.II]. Also included in Figure 9.6 is a graph showing thetemperature dependence of the AFC-77 cold-worked wire referred to in the precedingsection. It is seen that the wire retains considerable strength up to 1000T (540 C): thisadditional feature may be of great significance in fibre-reinforcement techniques.

107

r U.T-CONTAINING HIGH-STRENGTH STEELS

TEST TEMPERATURE. X

350 400 -VI 500 H C

800 900 «H1'EST 'EWPES4TLIRE. T

Fig. 9.?.— 100-hour creep-rupture strength of AFC-?". AFCOM). PyromeiX-15. UhraforUOl andseveral cobalt-free precipitation-hardening stainlesssteels. After E.J. Dixis ei al. [9.7]; data pointfor Ulirafort 401 after H. BRANDIS and A. VON DINSTHMN [9.6]; data for Pyromel X-15 after rVJ).

The 100-hour creep-rupture strength of AFC-77, AFC-260, PyrometX-i5 and severalcobalt-free precipitation-hardened steels is shown in Figure 9.7. AFC-77 is superior inthis respect to AFC-260 and Pyromet X-15, which are distinctly better than the cobalt-freesteels. According to the single data point available for Ultrafort 401, this steel is moreor less comparable to the PH steels. Typical creep and stress-rupture data for Pyromet X-15are given in Figure 9.8.

The favourable creep and stress-rupture properties of the various steels just consideredindicate that their thermal stability, or ability to retain adequate properties at high temper-ature, is good. However, most stainless maraging steels will show some thermal instabilitywhen held for prolonged periods in the 700 to 900 F (370 to 480cC) temperature range;the instability is manifested at room temperature by increased yield and tensile strengthsand decreased ductility and toughness. In this respect, a fair amount of work has beendevoted to AFC-77 in relation to its potential application in supersonic aircraft. Figure 9.9shows that holding the steel at 900:F (480C) increases its strength moderately. Theaccompanying decrease in toughness is not thought to be related to this increase but tobe due mainly to the removal of the last traces of retained austenite. No variation ofproperties was observed in the steel after prolonged exposure at 600" F (315 C). Table 9.4

p=T(ZWogt).I0-3fK.h]17

X 32 34p=T(20*tog 0.10-3 [Ra, hi

320

DO 2B0 Xt lEXPOSURE TIME hours

500

Fig. 9.8. — Creep and stress-rupture data forPyromet X-15 (Lar on-Mil!er plot). After [9.3\.

Fig. 9.9. — Effect of long-time exposure at 900"F(480°C) on the tensile properties and fracturetoughness of AFC-77. After D. WEBSTER [9.8].

Condition : 2000"F O095°C) - lh,O.Q. + 1100°F (595°C)-2+2h,A.C.

108

9 STAINi.ESS MARAGFNG STEELS — PROCESSING AND PROPERTIES

TABLE 9.4. — THERMAL STABILITY OF AFC-77.* After E.J. D L U S el al. [0.7]

Testtemperature

F

A) Mi

- 1 1 072

650

B) 10°

— 11072

650

C

U.

10'psi

r.s.,MN/nfi

cold reduction prior to ageing

22340

260237234

, cold reduction prior

—7922

340

29426425.

17921634

to ageing

202718201731

As-aged condition

0.2%

10-' psi

208182146

284255218

Y.S.,

WAV*.*

143412551007

195817581503

Elong.

111112

659

After !000-h exposure at 650 F (345"C)under 40,000 psi (275 MN/m-)

U.

10' psi

261215242

276262

T.S.,

MSInP-

185514821669

19031806

0.2";

103 psi

225202162

268230

Y.S.,

MS, m-

155i1393//.'/

18481586

Elong..

°'o

71010

h9

Condiliur, : 1900°F (1040°O- Ih. O.Q. / — 100cF (—73CC) - lh / 700°F (37OCC) - 2 +2h. A.C.

shows the effect of holding specimens of AFC-77 in two different conditions (aged, orcold worked 4- aged) for 1000 hours at 650°F (345°C) under a stress of 40.000 psi(275 MN/m2) on the tensile properties at —110°F (—80°C), room temperature and650°F [9.7]. The data confirm the pronounced strengthening effect of a 10% cold reductionprior to ageing, and show that long-time exposure at temperature under stress results inappreciable hardening, with little change in ductility.

9.2.3. Fatigue Behaviour

Rotating-beam fatigae tests have shown that the 10s cycle endurance limit of Ultrafort 401is 10.>,000 psi (735 MN/m2); this corresponds to a fatigue limit to ultimate tensile strengthratio of 0.43 [9.5], in good agreement with the value of 0.45 typically given as valid forhigh-sirength steels. It has been suggested that the fatigue limit of Ultrafort 401 can beimproved, a? is the case for non-stainless maraging steels, by decreasing the solutioningtemperature or by applying surface treatments such as nitriding or shot peening. The datalisted in Table 9.5 indicate that the fatigue strength of Ultrafort 403 is higher than thatof Ultrafort 401 [9.6].

The fatigue strength of AFC-77 bolts was recently shown to be greater than that of thehot-work die steel H-ll or the precipitation-hardening stainless steel PH 13-8Mo, and to

TABLE 9.5. — FATIGUE PROPERTIES OF ULTRAFORT 401 AND ULTRAFORT 403After H. BRANDIS and A. VON DEN STEINEN [9.6]

Steel

Ultrafort 401»

»Ultrafort 403

Specimen

smooi1!

notched (Kt =»

notched (Kt —

4.1)

4.O

CTmin

Omni

010

—10J

omax for rupture in

10-' psi

997443284733

107 cycles

685510295195325225

109

m i n i l iUMAIMNO HlUHSIRCNUtH STEELS

i ••• o Ml 1 ;Uvi ol ageing temperatureon :hc fatigue behaviour,.( \ ! C - " bolls. TEMPERING TCMPERATURCC

After IV W I B S I I R [».!•>]. a $Q ^f ^P - -vondr.ion : :iIW I III50 O - Ih. O.Q.: UNO F .f •. l i U d i »- H i . O O . . - HW I I - ' V C i - lh:aged lor 3 6 ; . . .AS MANUFACTURED 4 i-.!> at indicated Temperatures; thread rolled :- j• r -Nil | (4(H) Ci. Reiempcririi: for 2 hours at £ .! RETEMPEREO(60O"F-2h)KM K (.'15 c"i "as etTccwd to simulate high-tern- :--; "i AFTER THREAD ROLLING^

I'e-t .:onJiiion> : 1 4 in. dia. Polls, iuhjevted:o ienMon-tei;->ion fatigue tests vuith maximum and ;ii:nir.uiiiT load- of 4NK) j n J 4Milhs. rcs?ectiveK

be in excess of the minimum requirements for tvpical aerospace applic/.iins [9.14]. Theendurance limit of these bolls was also found to increase with increasing awing tempera-ture (Fig. 4.Kh.

^2 4. Lcrri'swn Resistance

The corrosion resistance of stainless maraging steels is structure as well as compositiondependent As regards composition, chromium has the greatest imnorianee, and numerousstudies on the Fe-Cr system have contributed to establish this fact. The role of cobaltappears to be of marginal significance. For instance, the corrosion resistance of Fe - 13Cr -10 and 20Co alloys in 10",, H;SO4 has been shown to be practically independent of theCo content [9.15]. This has been confirmed by potcntiokmetic studies in 25",', H2SO4.which have shown that an Fe-19Cr-l5Co alloy has practically the same behaviour as theferritic AISl 434 type steel of similar Cr content [9.11]. Molybdenum has a ?'.rong beneficialeffect on the corrosion resistance of annealed Fe-Cr-Co steels in 10% H2SO4 [9.15];copper is also beneficial, but its effect is much less marked when Mo is cc-piesent [9.16]. Thefavourable effect of these two elements as well as that of nickel is confirmed by the fact thatthe resistance of the Fe-19Cr-I5Co steel just referred to in 25°,, H2SO4 can be raised toor above the level of the austenitic AISi 304 steel through simultaneous additions ofl.4",,Mo, 1.7",,Cu and 1.3",,Ni.

Ageing of stainless maraging steels affect?, their corrosion resistance as a result of the struc-tural changes which they undergo. In particular, precipitation of Cr-rich phases such asR, 7, n. and (r^CrKMo results in decreased corrosion resistance through chromiumdepletion of the matrix. Generally speaking, the corrosion resistance is observed to decreasewith increasing ageing time and temperature [9.7/, 9.15, 9.17]. A compositional tiod-ification which enhances precipitation reactions will have the same effect. For instance,a 13.5Cr-5Co-6.7Ni-2Mo-0.5Ti steel was found to be unattacked both in the annealedand aged conditions, when exposed for 1000 hours to a 5% salt spray [9.18]; raising thecobalt content to 8 or 15 % did not modify the corrosion resistance of the steel in theannealed condition, but produced rusting when the steel was tested in the aged condition[9.19]. Similarly, the corrosion resistance of annealed 13Cr-10Co-2Moand 13Cr-20Co-2Moalloys in 10% H2SO4 was found to be the same; ageing at 930 or 1110°F (500 and 600cOhardly affected the corrosion behaviour of the 10%Co steel, but decreased that of the20 %Co steel.

Cobalt-containing stainless steels have been subjected to a variety of corrosive environ-ments, in order to compare their corrosion resistance with that of other stainless steels.Ultrafort 401 has been reported to to equivalent to AISI 410 [9.5], while Ultrafort 403in bar or sheet form exhibits no corrosion after exposure for 3600 hours in artificial seawater {DIN 50 900) or for 78 hours in salt spray (ASTM B 117) [9.6]. The corrosionresistance of Pyromet X-15 in the solution-treated condition is .stated to be at least compar-able to that of AISI 430 and possibly to approach that of AISI 304 [9.17]; ageing decreases

110

9 STAINLESS MARAGING STEELS — PROCESSING AND PROPERTIES

the steel's corrosion resistance. On the basis of 5°,, salt-spray test results, Pyromet X-23is claimed to be comparable to other precipitation-hardened high-strength steels of similaralloy content [9.20].


The stress-corrosion resistance of stainless maraging steels appears to exhibit a complexdependence on heat treatment. The discussion below will be mainly devoted to AFC-77,although some data are available for other steels [9.5, 9.6, 9.11, 9.20]. Results of stress-corrosion tests carried out on AFC-77 specimens aged at various temperatures are iistedin Table 9.6. Ageing in the 700-800 F (370-425 C) range results in good stress-corrosionresistance, while ageing at 900-F (480'C) and above leads to shorter times to rupture [9.13].The results plotted in Figure 9.11 were obtained under other experimental conditions, butare in qualitative agreement with those of Table 9.6. The stress-corrosion resistance passesthrough a maximum on ageing at 500 F (260 C). and through a minimum on ageingat 900 F (480 C): a marked recovery is observed on ageing at 1200 F (650 C). This variationis identical to that found for th- fracture toughness (see Chapter 8, Fig. 8.19), except thatthe A'|,. value does not show a recovery on ageing at high temperatures. A similar parallelismis observed between the K\,- and Klscc values : after ageing at 500F, AFC-77 has a KUecvalue of over 100,000 psi\ in. (110 MNm~3 '-): this decreases to about 50,000 psi\ in.(55MNm-3-) on ageing at 800 F (425C), and to about 10.000 psiy in. (11 MNra"-'-)on ageing at 1095°F (590'C) [9.22].

By analogy with the behaviour of 12°0Cr steels, it is believed that corrosion of AFC-77under stress involves a hydrogen embrittlement mechanism. The fact that overageingrestores the stress-corrosion resistance but not the toughness may be accounted for byassuming that, after ageing at 1200 F, the crack initiation period represents a major partof the total time to failure [9.21].

There is some evidence that retained austenite has a beneficial effect on the stress-corrosionresistance of AFC-77, as it has on its fracture toughness. Actually, fatigue crack growthrates in AFC-77 containing retained austenite were shown to increase only slightly onchanging the environment from dry air to 3.5%NaCl solution; in contrast, structures thatdid not contain retained austenite experienced a sharp increase in crack growth rate when

TABLE 9.6. — STRESS-CORROSION RESISTANCE OF AFC-77After E.J. DULIS and L. HABRAKEN [9.13]

Ageing

°F

700

750

800

900

1000

1100

ternperature,

"C

370

400

425

480

540

595

Hardness,

Re

49

50

51

51

52

52

Time

Specimen

>153

>153

.153

4

50

15

toda1

racture,*ys

Specimen 2

>153

>153

>153

5

84

>153

* Test conditions : two specimens 0.05 in. (1.27 mm) thick were aged for 2 -t 2 hoursat indicated temperatures (initial condition : I900°F (I040°C) - 30 min, O.Q.T refrigeration), then loaded as ben| beams to produce a stress of 124.500 psi(860 MN/m-) and exposed in air while subjected to a daily aspersion with a 5% NaOsolution containing a wetting agent. pOr comparison, cobalt-free stainless steeiswere tested under the same conditions: the results were as follows :AM .150 : 14 and 15 days AM 355 : 5 days PH I5-7MO : > 153 days.

100

80

e|j-eo

' 2C —

1'

TEMPERING TEMPERATURE, °C

2DD 300 a)0 500

I

o

111 • • J

rt

i . i

/

yA' \

a

1

\

\/ \. O FAILURE CAUSED BY PROPASA- \

TiuN OF THE FATlGLfc CfiftL* '• • FAILED OUTSIDE THE PLANE OF

FATIGUE CRACK DUE TO EDGECRACKING

I I

i

< /

Vn

1

/

400 BOO BOOTEMPERING TEMPERATURE, T

1000

Fie. 9.11. — Effect of ageing temperature on siress-corros-ion resistance of AFC-77. After R.T. AULT el al. [9.21].

Centre-notched fatigue-cracked tensile specimens loaded at 90%of notch tensile strength ani exposed to acidified 3.5%NaClsolution (pH 1.5). ConJition : 19000F (1040°C) - lh, O.Q.;—100°F (—73°C)-30min; aged for 2+2h at indicated temperatures.

I l l

( OHM r-l/ONTMMNLi HU.ll STRENGTH STUE1.S

I \BI I ' )" . SI RISS-COKROSION HI SIS! A M I (» : L I I KA1OR1 40JIN ALk.-ULU 5',, Nat I. After H. BRA.NDIS and A. VON o t s STEINI.N [V.fi]

Specimen *

Smooth

Notched (K: 11

Smooth

10 'p

242

242

Applies stress

Te\ts iir r

Tests HI

W.V m-

tif/n wrnptrat il'C

:i; F : inn c

I4S0

Stress

Y.S.

0.900.9?0.900.95

0.700.90

Timeto fracture,

days

I1?.300

-42•45

13, 14

• RL-niovcJ from -1 4 ir.. iW mm> bar.

the same change was made [^.--]. Further siudies have shown that maximum stress-corros-ion resistance in AFC-77 is obtained using relatively low austenizing temperatures asso-ciated with low ageing temperatures. On the other hand, a 20"„ cold reduction afteraustenizing at 1900 F (1040 C) has a detrimental effect, even when followed by ageing at700 to 800'F (370 to 425 O : however, a 10",, cold reduction applied after austenizing at1800 F (980 C) does not adversely affect tne stress-corrosion resistance [9.7].

Stress-corrosion data for Ulirafort 403 are presented in Table 9.7. Although the lest condi-tions are not identical with those described in Table 9.6, both sets of results are indicativeof the good stress-corrosion behaviour of the steels under consideration, at least for thetreatments involved. It should be added, however, that a recent evaluation of massiveforgings made from AFC-77 and " Alloy B " has not shown their stress-corrosionresistance to be higher than that of high-strength medium-alioy steels and other high-strength stainless steels [9.23]. Since several of the latter have also been reported toexhibit severely degraded flow tolerance in a corrosive environment as the section sizeis increased, it would appear that the problem of low KUcr values (10,000 to 20,000psi \ in., i.e., 11 to 22 MNm ' -) experienced with the AFC-77 and "" Alloy B " forgingsmay be a characteristic of high-strength stainless steels in general.

TIME, hours

Fig. 9.12. — Oxidation behaviour of AFC-77and several cobalt-free high-strength stainlesssteels exposed in still air at 120CTF (65O=C).After E.J. DULIS and L. HABRAKEN [9.13].

112

9 STAINLESS MARAGING STF.ELS — PROCESSING AND PROPERTIES

4.2.6. High-Temperature Oxidation Resistance

For steels such as AFC-77, which are intended for use at both room and high temperatures,the oxidation resistance is a significant characteristic. The weight gain curves recordedon various steels, including AFC-77, exposed in still air at 1200 F (650 C) are shown inFigure 9.12; the oxidation resistance of AFC-77 appears io be superior to that of the othersteels selected for comparison [9.13].

9.3 Secondary Processing

Cold working of stainless managing steels is best performed n the solution-annealedcondition for the low-carbon grades, and in the quenched - aged condition for the carbon-containing grades. For instance, a double ageing treatment of 2 — 2 hours at 500 F(260 C) is recommended before cold working AFC-77 [9.1].

Machin3bility, which is reported to be similar to that of other maraging steels, is best inthe solution-annealed condition for the low-carbon grades (e.g. Pyromet X-15, PyromstX-23, Ultrafort 401) or in the annealed condition (2 — 2 hours at l-*00 F, i.e.. 760 C)for AFC-77 [9./. 9.3 to 9.5]. In these conditions, the steels have relatively low work-hardening rates, which favours machinability. Because of their high ductility, siecls sucha; Pyromet X-23 can also be machined in the aged condition; however, lower speeds andmore rigid work supports should then be used.

Surface treatments ich as nitriding increase the wear resistance, and probably also thefatigue strength, of stainless maraging steels [9.5]. Gas nitriding Ultrafort 401 at 890 F(475C) for 40 hours produces a 4 mil (0.1 mm) thick surface zone with an HV0.5 hardnessof 1340. The nitriding conditions coincide with those for ageing, so that only one treatmentis required.

Welding, generally in the solution-annealed condition, can be satisfactorily performed bythe TIG process with or without a filler metal of matching composition. After welding, asimple ageing treatment is sufficient to achieve a joint strength similar to that of the basemetal; however, for optimum strength and ductility (particularly in heavier thicknesses),a post-weld solution treatment is usually recommended. TIG welding of Pyromet X-15strip has resulted in joint efficiencies greater than 90% at maximum strength levels [9.3];similarly, joint efficiencies of 95% have been reported for AFC-77 [9./]. In the case ofthe latter steel, it has also been shown that no susceptibility to hot cracking developsduring welding; however, due to low notch toughness in the heat-affected zone afterageing, the joint should be carefully designed [9.24].

9.4. Applications

Cobalt-containing stainless maraging steels have been designed for applications requiringhigh strength and good oxidation and corrosion resistance up to about 1050°F (570°C)for low-carbon grades, and up to 1200°F (650°C) for higher-carbon ones. Potential appli-cations include steam turbine and jet engine components (e.g. blades and compressor discs),aircraft structural parts, gun barrels and gun-barrel inserts, bolts, fasteners and springsoperating at elevated temperatures, and extrusion, forging and die-casting dies. Althoughmany of these applications are still under evaluation, it appears that significant aâge maywell develop for steels like AFC-77 in the aerospace field as fasteners [9.14] or as structuralcomponents with the advent of Mach 3 supersonic aircraft [9.25]. AFC-77 has also beenextensively evaluated as a stainless bolt material for operation in corrosive environments[9.14]. Similarly, the suitability of Ultrafort 401 as a pressure-vessel material for thechemical industry has recently been demonstrated [9.26].

113

(.HUM l-l l A I \ i \ I M . lilt.II-SIR1 Nt . IH S T t f l S

ill. tONCl. l SIGNS

Cobalt has been shown to play an important role in the three families of high-strength>ieel> reviewed in the present volume, ri:. carbide-strengthened, Ni-Co-Mo maraging, and-lainless maraging. Very r'iefly. in addition to exerting a small solid-solution hardeningetiec, cobalt increases the \/< temperature: this factor has been used to advantage inuiHim ractuig the depressing cttect of the other alloying elements present in tH~se steels.Cobalt llso retines the martensitic structure and leads to retention of the dislocation sub-structure at higher tempering temperatures, thereby giving rise to a finer precipitationof dislocation-nucleated carbides or intermetallic compounds. Final'.', cobalt may affectthe precipitation reactions themselves, e.g. by decreasing the solubility of molybdenumand chromium, and may induce ordering in the matrix.

\> a result of these efleets particularly the obtention of a tine dispersion of the strength-ening phases, the steels under consideration exhibit attractive combinations of strengthand Toughness. The Ai, >'.v. yield strength relationships shown ia Figure 10.1 recapitulatethe information presented in Chapters 4, 6 and ''; they clearly confirm the superiority ofthe cobalt-containing grades over other high-strength steels, except perhaps the TRIP(transformation-induced plasticity) steels, for which A'i,- value1 as high as 130,000 psiy in.1140 MNm ; -) at yield strength levels of 250,000 psi (1725 MN/m-) have been predicted[10.1]. Specifically, the diagram illustrates the exceptional position of the 10Ni-8Co-Cr-Mosteel and the favourable one of the HP 9-4-20 and HP 9-4-30 grades at the 180,000 to

,1000YIELD STRENGTH. MN/m !

1600 VMO 20C0 2200

160 200 240 2S0CIELD STRENGTH, 10> psi

320

Fig. 10.1. — Strength/toughness relationship for various categoriesof high-strength alloys, showing position of cobalt-containing steels.References as in Figures 4.7, 6.1 and 9.1, except for TRIP steelsand ,'J-Ti alloys, after A.R. ROSENFIELD and A.J. MCEVILY [10.1].

High values of Klc may not always have been determined accord-ing to ASTM specifications, but trends shown should be significant.

114

ID. CONCLUSIONS

200,000 psi (1250 to 1380 MN/m-) yield strength level, as well as the unique combinationof properties afforded by the Ni-Co-Mo maraging steels over the 200,000 to 375,000 psi(1380 to 2580 MN/m2) range. Most of the cobalt-containing stainless maraging gradesalso offer good to outstanding compromises.

However, the above combinations of strength and toughness are not the only favourablefeature of the cobalt-containing high-strength steels. Their other properties lie at leastat the normal levels for steels of this class, and in many cases are well above them.Particularly worthy of note are the very high mechanical properties of the 18Ni-Co-Mosteels at cryogenic temperatures, and the superior stress-corrosion resistance of theHP 9-4-20 and 10Ni-8Co-Cr-Mo steels. Some properties pertaining to individual steelsalso deserve a special mention. For instance, the combination of good high-temperaturestrength and high magnetic induction of the 15Ni-15Co maraging steel is expected to leadto its application in high-speed electrical motors and generators: similarly, the high hardness,strength and temperature capability of the 18Ni (350) and l3Ni (400) grades should favourtheir use in tooling and die applications at moderately high temperatures, while the com-bination of high fatigue strength and stress-corrosion resistance of AFC-77 pleads in favourof its use for high-strength bolts.

It has been shown in the preceding chapters that the steels under review do not present anyproblems as regards processing. In particular, their hot workability, cold formability.machinability and weldability are usually more than adequate. Here again, some speciallyattractive features of specific steels are worth pointing out. These include the excellentweldability of the HP 9-4-20, 10Ni-8Co-Cr-Mo and I8Ni-Co-Mo grades, and the fact thatmaraging steel parts can be machined to their final dimensions in the annealed condi-tion, since the subsequent ageing treatment causes minimal dimensional changes anddistortion. Appropriate techniques have been developed to fabricate the steels, not onlyby melting and casting followed by hot and cold working, but also by powder metallurgymethods. Some grades can also be used as castings; this applies, in particular, to a speciall8Ni-Co-Mo grade tailored Tor increased fluidity and castability.

As regards present and potential uses of cobalt-containing high-strength steels, theyare currently restricted to three major fields : (I) aerospace, aeronautic and maiine appli-cations (rocket motor cases, ianding-gear links and other highly-stressed aircraft compo-nents, hulls of deep-sea vessels, etc.); (2) structural and machine components (pressurevessels and reactors, bolts and fasteners, bearings, springs); (3) tooling applications (die-casting dies or components thereof, extrusion rams and dies, cold-forming dies, camfollowers, tooling fixtures, etc.). On account of their adequate strength retention atmoderately high temperatures and their good oxidation and corrosion resistance, thecobalt-containing stainless maraging grades are also being considered for use in steamturbines and jet engines, e.g. as blades and compressor discs. It is difficult to predict theways in which the cobalt-containing high-strength steels will be used in future, except inthe general terms employed in Chapter 1 (see Fig. I.I). However, it can reasonably beexpected that the rapid technological advances which characterize our time will make gooduse of the superior, and sometimes exceptional, combinations of properties possessedby these steels. For instance, there are at present strong indications that the high speeds(50,000 to 100,000 rpm) necessary for the ultra-centrifugal enrichment of uranium willrequire the use of the higher-strength maraging grades, and this may eventually lead tovery significant consumption of these steels.

115

( OH \l I-I. ONI A1NING HKIII-STRI NCMII S I E H S

REFERENCES

1. INTROIHCTION

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[2.24\ Yi«. R.B.G.. The Effects of Some Alloying I lemenls on theTransformation of Fe-22.5'\,Ni Al!<i>s. Tram. A.I.M.I-:.. 227.S,S4 (1963).

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[2.31\ SPHCH. G.R.. and SWANN. P.R., Yield Strength and Trans-formation Substructure of Quenched Iron-Nickel Alloys.J. Iron Steel lust., 203. 480 (19651

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[J.4ftJ FLORIKS. S., The Physical Metallurgy of Maraging Steels,Met. Reviews. 13, 115 (1968)

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3. CARBIDE-STRENGTHENED STiiELS —PHYSICAL METALLURGY

i.'./J PLRRY, T.b., PooLli, S.W.. and MAIAS, S.J., Development ap.dMelting of an Ultrahigh Strength 9"oNi-4"uCo Steel, in"'Electric Furnace Proceedings", A.I.ME., New York,Vol. 20, 1962, p. 308.

[3.2] PASCOVER, J.S., and MATAS, S.J., •• Relationships BetweenStructure and Properties in the 9Ni-4Co Alloy System",A.S.T.M. Spec. Tech. '" jbl. 370, 1965. p. 30.

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[3.4] KONKOL, P.J., RATHBONE, A.M..and GROSS, J.H.. Developmentof 170/200 ksi Yield Strength Ni-Cr-Mo-Co Weld Metals forConstructional Steels, Welding J., / to . SuppL. 45, S25i U9M».

[3.5] DABKOWSKI, D.S., KONKOL, P.J., PORTI R. L.F.. and KATHBONL,A.M., Nickel-Cobalt-Chromium Steel, U.S. Pat. 3.502,462,24 March 1970 (Appl. 29 Nov. 1965).

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[J.7; BIRKLE, A.J., and PORTER, L.F., The Effect of Cobalt on theStrength and Toughness of Ni-Cr-Mo High Y'ield StrengthSteels, United States Steel Applied Research Laboratory Rept.on Project No. 40.018-002(16) (AD 600 790), April 1964.

[J.u] STONESIFER, F.R., and SMITH, H.L., Characterization of GTAWeldmems in 10Ni-8Co-2Cr-lMo Stee', Na.al ResearchLaboratory Memo Repl. 2466 (AD 7<>o 111), June 1972.

[3.9] SPEICH, G.R., DABKOWSKI, D.S., ard PORTER, L.F., Strengthand Toughness of : c-lONi Allo;s Containing Carbon,Chromium, Molybdenum and Cobalt, Met. Tram., 4, 303 .(1973).

[3.10] HOTZLER, R.K., MACIAG, R.J., WOSKCWIAK, A.J., andAPRIGUANO, L.F., Polytechnic Institute of Brooklyn, Brooklyn,N.Y. " Precipitation Characteristics of a 10"oNi Alloy SteeP",Paper presented at 5th Annual Spring Meeting of A.I.M.b.,Philadelphia, Pa., May 29 - June 1, 1973.

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[3.12] AULT, R.T., HP 9-4-20 Steel Data. Republic Steel Corp. Tech.Rept. TR 120IS-132. March 11, 1968.

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K V I I M I . I ' K; i i \ . s -\ . .ui!.t II raciutc Ioiie.hm.-v. ol l 'N>-4l .I ci-ipcr ins: .iiul Xii-foniiing. '

mis. M . \ lo\» Strength jndSteel JS Allected f> Slum

• / . * / > / ' i f < . ' " • " ' 7 ' • * ' . * 4

.'' .'•>; I i ' \ . I H . .ni.t R n s f s . s u IN. A H . I hi- Propel tics a n d Mivlo-stiih.li ie ol I O N I - 2 1 r - l M o -si o S l ed . J Mai,' . ft. " V s t l s T l i

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. . .••-: I KI s. u i s . i . VI . I o m n b u t i o n to the Study of the HealI tc i i ' i icni ami Mechanical Piopeit ies **< s>Ni-4i o Steels. He*.\'.':••':./,«. V . / . ' • « . .S. '«• W . v . l l . l»sl i |sii.M.

.•' . v: I i si i> U . Sum R. R I . R i m IM k. S Ci . and ( > K I i v S J .1'i.iMis I i iw i" Binary Subsi i tui ional Alloys ol H I C Iron1 ikxis ol s t i j i n R a t e . I c m p e r a i u r i ' a n d Alloy C o n t c n l . hart*1 ,\ 1/ . h2. iÔ i |'»t*»i

• _'•': I I K . I I M V . I I I , 1 |> . [ i si A. A R . a n d H u i . A M . I he >»Ni-4l o

Steel-. H M H M e m o 22l>. IJclobcl l'-'M'

• ' , \ : l i " W > v S . I hi Propert ies of l o w - C a r b o n Iion-Nu.ki: l-V l . iommni Mar icns i ics . Iran* I I \l t . 2.16. 142^ i I '"- )

;.•. '-", i nvNimiik.. V K . H I K I H . J P . and D i n s . I J . . I.tlect otv . ' ,!l on l e n i p e n n g Tool and Alloy Steels. Jiun* ISM.

;.'.. '. ') vi!AM»u>k. \ . K . HtRlH. J R . and Di 1 is, I . . ! . I Ifect ofCobalt on C a r b o n Activity and OilfuMvuv in Steel . Iran*.

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Z).

4. (WRBIDK-STRKNGTHENtD STl-.lXSPROC KSSING AND PROPtRTlt.S

[4.1 \ CiK.viNi\tt.i>. T.P.. I 1st A. A .R.anJ HAIL. A.M.. Fhe9Ni-4CoSteels. DMIC Memo 220. October Whf..

[4.2\ Putm. I I .. Pi«>ir. S.W.. and MAT vs. S.J.. Development andMelting of .in Lltrahigh Strength 9",,Ni-4",,Co Steel, in" t.lectnc Hii'iace Proceedings " . A I. M.I... New York. Vol. 20,WfO. p. <(l.v

[4.J] MAI vs. S.J.. Influence of Impurities am! Related ttlecu, onStrength and Toughness of High Strength Steels. Metals Eng.Qmiri.. 4 (2). 4S (19f.4|.

[4.4] BIRKLK A.J.. and PORTER. L.F.. The Effect of Cobalt on theStrength and Toughness of Ni-Cr-Mo High Yieid StrengthSteels. United States Steel Applied Research Laboratory Rept.on Project No. 40.018-002(16; (AD-600 790), April 1964.

[4.5] RATHBOM . A.M.. Forum : Where We Stand in Welding High-Strength Steels. Where Two New Grades Fit into the Picture,Metal Progr.. 95 (2). 67 {1969).

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[4.7] - Properties of Republic HP 9Ni-4Co Steels " . Republic SteelCorp.. Cleveland. Ohio, !96h; see also Alloy Digests SA-179(July 19651. SA-27X (July 1972). SA-IK1 (Aug. 1965). SA-270(Dec. 1971). Engineering Alloys Digest Inc.. Upper MontclairN.J.

[4.H] " Properties of HP 9-4-20 Steel " , Republic Ste.-I Corp. DataCompilation, September 1972.

[4.9] STOSESIFER. F.R., and SMITH, H L., Characterization of GTAWeldments in 10Ni-8Co-2Cr-lMo Sleel. Naval ResearchLaboratory Rept. 2466 (AD-746 111), June 1972.

\J ID]

\4 , ' / ]

(.•KOI M M i u. I P . I niiiii n ( I . I IIIII . R I . and SIMONS.R v . High'SiicniMli Mi-el ' i \ i - i i o. DMK Pioicssts andPropeities Handbook trtw, I olumbus. ohm

Si ' i t tM. Ci R . l>Aiik<>uski. U S a n d P O K I I H . 1 I . Si iengll iand l o u g h n e v , of l e -HINi Alloys Conta in ing l . ! i b n n .I iMoniium. MoKbdi-iuini and C o b a l t . \h I Iiunt. 4. M\\ll>»"3i

i t M siNn. H P . ICIIMIL- Siii-nglh of I ighl C on\'.\ mt io ,Steels Between ~l> and 320 I . J Mum . 4. -l\< (I'li.-Ji

CiK.iv,. J II , I S Sleel C o i p . Pr ivate t .

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At i I. R I . H P '1-4-20 Steel DaiRept ! R 12 •"

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Sli'i'l *. o i p lci.h

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[4. IV\

Repl !R I2(li.s-I;. M.IKII II. r'l.h.

Pwi .Ufc . J S . PiopeiltfM'l 1'ioduUion Product of lll"'-4-31).Republic Steel Re-eauh C en'er Pubhcn . Dvcembci S. | W

P A M " \ I « . 3 S . and M \ n s . S J . Relationships lleivveenMim-ture and Piopeitics in the l>Ni-4l o Alloy System, in••Striiituie and Piopcities of I Iiiahijih-Sticnglh Steels".\ S I M Spec lech Publ No .'"(>. l » v p W.

K.\usti. I).. Ki u s . S.A., and CIHIIN. VI., l-'km Strength andFracture Toughness oi 9Ni-4Co Steel as AlVected by Strainlempenni! and Ausformmg, Mclat\ lint:. Quart.. 7 (4), 54(196-).

14.

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14.24]

K A I ISH. [ l . . O n u s . M. . and Ki u v S A.. Strain f t n i p e r m g ofBatnite in yNi-4C'ivO.45C" Steel. J. Muter.. 5 . KM (19701.

\ ' i s n \ T v s k ^ . C . and S I I I V I I K V > A I I ) , L A , Plane Strain i r a c l u r eToughness at Room and Sub/em Temperatures, in " IracuircToughness Testing a( Cryogenic Temperatures " , A.S.T.MSpec. lech. Publ. 49d. i97 l , p. 3.

" Mechanical-Property Data on HP 9Ni-M.o-0.20C Steel"Balletic Memorial insiituie. Columbus, Ohio. September 19"2.

DARSOM. J.M.. IMIIOI. I J . and R<u H . ST. . I atigue-CrackPropagation in Hiuh Yiclu Streni'th Steels, f.iii'. riimiirfMechanic*. 2, 301 (1971).

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C ARTHI. C.S.. Crack Kxtension in Several High-Strength SteelsLoaded in 3.5 "„ Sodium Chloride Solution, Boeing Co. Rept.No. D6-I97Ô, November 1967.

STONESIFEU, F.R.. SMITH. H.L., and ROMINL, H.E. Propertiesof Hot-Pressed lONi-Cr-Mo-Co Steel. Naval Research Labor-atory Memo Rept. 2065 (AD 699 530), November 1969.

[4.251 Mt'NOEB, FI.P., High-Strength 9",,Ni-4",,Co Steels and TheirApplications. Cahnll. No. 44, 127 (1969).

[4.26] Rits. G.D.. and POOLF., S.W.. Welding of Quenched amiTempered 9Ni-4Co Steels, Welding J.. Res. Suppl., 45- 465s(1966).

[4.27] COLTSOLRADIS, D., LAMBERT, N., DRAPIFK, J.M., andHABRAKFN, L., Microstruclure and Properties of 9Ni-^CoHigh-Strength Steels, Cobalt. No. 37. 192 (1967).

[4.28] SAVAS. J.. BECKER, P.C.. and MATAS. S.J., Development of a1X0 ksi Yield Strength Stress-Relievable 9Ni-4Co Filler MetalComposition. Weldinp J., Rex. Suppl., 48, 479i (1969).

[4.29] UCHIDA, J.M., Welding Evaluation of New High-StrengthQuench-and-Temper Steels, Welding J., Res. Suppl., 45, 31 j(1966).

[4.30] RATHBONE, A.M., Review of Recent Armour Plate Develop-ments, Blast Furnace SWei Plant, 56, 575 (1968).

[4.31] MATAS, S.J., Republic Steel Corp., Private communication,1973.

118

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5. Ni-Co-Mo MAR AGING STKKLSPHYSICAL MKTAl.lXR(;Y

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[5.11] LKiiMiKi. P.. and I_"ASTAI;NI'. J.L.. The Contribution of Elec-tron Microscop> to the Slud> of the Technological Propeniesof Maraging Steels. Met. Constr. Mec. 100. 105 (I9M5).

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[5.17] PATTERSON. W.R., and RICHARDSON, L.S.. The Partial Substi-tution of Manganese for Nickel in Maraging Steel. Trans.A.S.M.. 59, 71 (1966).

[5.18] TcrrNELL. G.W., and CAIRNS. R.L.. 18°,, Nickel (350) Mar-aging Steel, Tram. A.S.M.. 61, 798 (1968).

[5.19] BOESCH. W.J., and COWAN. T.W. (Special Metals Corp., New-Hartford, N.Y.), Evolution of a Commercial 400 ksi GrtdeMaraging Steel, Paper presented at A.S.M. Materials Engineer-ing Congress, Detroit. Mich., Oct. 14-17. 1968

[5.20] MIHALISIN. J.R.. and BIEBER, C.G.. Theop-'ical Strength withIron-Nickel Maraging Steels. J. Metals. 18. 1033 (1966).

[5.2l\ FLOREEN, S., A Study of Maraging Steels with Higher Cobaltand Reduced Molybdenum and Titanium Contents, J. IronSteel /;«;., 207, 484 (1969).

[5.22] KAUFMAN, L.. and COHEN, M., The Martensitic Transformationin the Iron-Nickel System, / . Metals, 8, 1393 (1956).

15.2.'] MIIIAIISI"., JR.. F fleet of Carbon Content on TransformationStructure of Iron-22",, Nickel Alloys, in "Fifty Years ofProgress in Metallugniphic Techniques". A.ST.M. SpecTech. Puhl. No. 430. I9ISH. p. 250.

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[5.21] SWANSON. W.D.. and PARR. J.G.. Transformations in Iron-Nickel Alloys. J. Iron Steel Insl.. 202. 104 (1964).

[5.27] OVVI.N, W.S.. and WILSOV. FA.. A Note on Mas-,i\e Struclure.in '• Physical Properties of Martensite and Bamiir ". Spec.Rept. No. 93. Iron and Steel Institute. London. 1965. p. 53.

[5.2H] BOI R<,K1T. J.. MAITRfflLRRf. P.. M»\>« . J.. and TH(IM»S. B.Hardening Fe-Ni-Mo and Fe-Ni-Co-Mo Martensitic Alloj- h>Tempering. IRSID Internal P ept. No. lift. August 1971.Part I: Electron Microscope Study of Coherent Precipitationin Fe-Ni-Co-Mo Maraging Alloys. Paper presented at5th International Materials Symposium. Universitv of Califor-nia. Berkeley. Calif.. September 13-17. 1971.

[5.2V] MILLER, G.P., and MITCHELL. W.I., Structure of Nickel-Cobalt-Molybdenum Maraging Steels in the Air-Cooled Condi-tion. J. Iron Steel /its/.. 203. 895 (1965).

[5JO] BONIZEVVSKI. T.. Hydrogen-Induced Delayed Cracking inMaraging Steel. Brit. Weld. J.. 12. 557 (19ft5l.

[5 31] Yf.o. R.B.G., The Effects of Some Alloying Elements on theTransformation of Fe-22.5°oNi Alloys. Trans. A.I.ME. 111.884 (1963).

[5.32] HAMMOND, CM., The Development of Maragina StainlessSteels Containing Cobalt. Cobalt. No. 25, 195 (19641.

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[5.3-1] SHYNE, J.C.. ZACKAY. V.F., and SCHMATZ. D.J.. The Strengthof Martensite Formed from Cold-Worked Austenite. Trans.A.S.M., 52, 346 (1960).

[5.35] WARD, D.M., MILLER, G.P., and BRIUGES, P.J.. Effect ofMartensitic Transformation on Cold Formabilitv of MaragingSteels, Rev. Met., 68, 107 (1971).

[5..?rJl CHILTON, J.M., and BARTON. C.J., Identification of Strength-ening Precipitates in 18Ni(250) Aluminium, Vanadium audTitanium Maraging Steels, Trans. A.S.M.. 60. 528 (W67).

[5.37] MARCUS, H., SCHWARTZ, L.H.. and FINE. M.E., A Study ofPrecipitation in Stainless and Maraging Steels Using theMossbauer Effect, Trans. A.S.M.. 59, 468 (1966).

[5.3S] REISDORF. B.G., Identification of Precipitates in 18, 20 and25"O Nickel Maraging Steels, Trans. A.S.M., 56, 783 (1963).

[5.39] BAKER, A.J., and SWANN. P.R., The Hardening Mechanism inMaraging Steel. Trans. A.S.M., 57, !008 (1964).

[5.40] BANDI, W.R., L.UTZ, J.L., and MELNICK. L.M., The Extraction,Identification and Quantitative Delermination of Second PhaseCompounds in Aged 18°i Nickel Maraging Steels. J. IronSteel Inst., 207, 348 (1969).

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, ; -'•; ii\hU'»M>. K P . -UK! JUMM. K.I).. 1 lection Micioscopc Mud>of the structure .>t a -S",.Ni. Ti. AI Maraging Steel. V. taMi , ; ;>isr . JIM. <I2 I I * I M .

|> J J | PKM-IIK. J M . \ m m R. I'.. M u . M i . A . and Cm Tsol RUMS,\i . Mr Kiui.il Suidv of 4lH>-liradc Maragmg Steel. Cohalt.\ i - <o. ;-> 11'<"11-

| \ J . ; | H I V K I I . H R . anJ M M SIR. J ' . . Transformation Kineticsin Mar.iiitf-I'vpf Ke-ls".Ni Sleels. n •• transformation andM.irdeiubilny in Steels " . Climax Mikhdenum Co.. AnnVibor. Mich.. i^h.s. p. I.-..V

|5-'ftj I AMHIKI. \ . VlMull l . P.. DkAI'llR.J M . and CtH TSt>t RADIS.II . Hardening Mechanism- in Manuring Steels. Mem. Set. fii'i.\h: . 57. 4 . ^ ' t W l l i .

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[.'..'-I MILLFR. G.P.. and MITCHFIL. W.I.. Structuie and HardeningMechanisms of IS",, Nickel-Cobalt-Molybdenum MaraeingSteels. J Iron Sieel / « . . 203. SW (l%5l.

[5..V1 Cbi'TSOi v \uis. D . DKM'JFH. J.M.. DIDFRRK FI. E.. and HAORA-Kts. L.. I'recipitution Hardening in High-Streneth StainlessSteels. Cohali. No. ib. 144 (l%"t.

[?.54] PATTIRSDN. R.L.. and VVAVMAN. CM.. The Crystallographyand Grov.'.h of Panial|\-Twinned Martensite Plates in Fe-N'iAlloys. Ada Met , 14. 34"* (19661.

[5.55} MIHALISIS. J.R.. Age Hardening and Structure of Iron-Nickel-C'oball Alloys. I ram. ASM.. 59, CO (19661.

IS.56] FLOREEN. S.. and DtfKtR. R.F., Heat Treatment of 18"oNiMaraging Steel. Trans. A.S.M.. 55, 518 (1962).

[J.J7] SI'OOSTER. S., RACK. H.J.. and KAUSH, D., Neutron DiffractionAnalysis of Atomic Arrangements in a Maraging Steel. Met.Trans.. 2, 2306 (1971).

[5.5S] RACK. H.J., and KAUSH, D., The Strength and FractureToughness of 18Ni (350) Maraging Steel. Met. Trans., 2, 3011(1971).

15.59] FLOREEN, S., and SPEICH. G.R.. Some Observations on theStrength and Toughness of Maraging Steels, Traits. A.S.M.,57, 714 (1964).

[5.60] PETERS, D.T., and CLPP. C.R.. The Kinetics of Ageing Reac-tions in I8"oNi Maraging Steels. Tram. A.I.ME 236 1420(1966).

[5.61] BANERJEE. B.R., HAUSER, J.J., and CAPENOS, J.M., Role ofCobalt in the Marage Type Alloy Matrix, Metal Sci. J. 2 76(1968).

[5.621 MINER, R.E., JACKSON, J.K., and GIUBONS, D.F.. InternalFriction in !8°oNi Maraging Steels, Trans. A.I.M.E 2361565 (1966).

[5.63] DETERT, K., Investigation of the Precipitation Behaviour ofHigh-Strength Maraging Steels, Archiv Eisenhiiitemv., 37,579 (1966).

15.(W| IdKtNAi.A. Y., and OKITA. T. (Kyushu University). Paperpresented at 6Sth Spring Meeting of Japan Institute of Metals.I97|.

[5.65] PITERS, O.T.. and FIIIRHN. S . Precipitation Hardening ofFerrite and Martensite in an I c-Ni-Mo Alloy. Tram. A.I.M.E.,245. 2021 (IWJi.

[,5.601 MARCIS. H.I , PUSHUP. J.N . and FINI. M.I... Precipitationin 17-7 I'M Swmless Sieel, Trans ASM.. 58. 176 (l'J65j.

(5.67) PETIRS. D.T., Precipitate Reversion in IS",,Ni-Co-Mo Steels.Trans. A.I. V/.f... 239. I9S1 (19h7i.

[5.68] HoRNTKXiiN. E.. Clustering in an i Iron-Molybdenum SolidSolution. J. Appl Plus , 32. 135 (iy6l).

[5.6VJ MARCIS, H.. F I M . M.I'... and SIIIWART/. I II., Mossbauer-Lffect Study of Solid-Solution and Precipitated f-e-RichFe-Mo Alloys, / Appl. Pliys., 38. 4750 (I967i.

[5 Tt'] ERR-SSOS. T., MUI Rtkis. S., and Cunts. J.B.. Clustering inIron-Molybdenum Allovs. J Mater. Sci., 5. 901 (1970).

[5.'/I K.iM)SitirA, T.. TDKINAIIA. V , and TOVIISMIMA, T.. Calon-metric Analysis of I.o»- temperature Maraging Stainless SteeU,Mppon Kinzoku Gakkai-Si, 33. 260 (1969).

[5.7J] TtiKLNAiiA, Y. and KINOSHITA, T.. Resistometric Analysisof Low-Temperature Ageing in Maraging Stainless Steels,Trans. Japan Insl. Metals, 12, 250 (1971).

[5J3] PRIESTER. L., Polentiokinetic Study of Maraging-Type Alloys,Mem. Sci. Rev. Mel., 67. 707 (1970).

[5.741 PETERS, D.T., A Study of Austenite Reversion During Ageingof Maraging Steels, Trans. A.S.M., 61, 62 (1968).

[5.75] FLOREEN, S., and DECKER. R.F., Maraging Steel for I000FService, Tram. A.S.M., 56, 403 (1963).

[5,761 Bui, NAM, and DABOSI. F., Contribution to the Study of theEffect of Molybdenum on the Ageing Kinetics of MaragingSteels. Cobalt, No. 57, 192 (1972).

[5.77] SERVANT. C , and CIZERON, G., Effect of Cobalt and Molyb-denum Additions on Structural Transformations in Mar-aging-Type Alloys, Mem. Sci. Rev. Met., 66. 531 (1969).

[5.75] ViALATTE, B-. and DUBOIS, B., Investigation of the DilatometricBehaviour of Maraging Steel 18Ni-8Co-4Mo, J. Iron SreelInst., 209. 147 (1971).

[5.79] LEGCNTJRE, P., Some Properties of Maraging-Type Steels,Cobalt; No. 29, 171 (1965).

[5.80] THEVENIN, J.-p., CIZERON, G., and LACOMBE, P., AusteniteStabilization in 18°;Ni Maraging Steels by Interrupted Quench-ing, Mem. Sci. Re . Met., 68, 75 (1971).

[5.5/] KESSLER, H., and PITSCH, W., On the Nature of the Martensiteto Austenite Reverse Transformation, Acta Met., 15, 401 (1967).

[5.82] JANA, S., and WAYMAN, CM., Martensite to F.C.C. ReverseTransformation in an Fe-Ni Alloy. Trans. A.I.M.E., 239,1187(1967).

[5.83] SHAPIRO, G., and KRAUSS, G., The Crystallography of theReverse Martensitic Transformation in an Iron-Nickel Alloy,Trans. A.I.M.E., 239, 1408 (1967).

[5.84! GOLDBERG, A., and O'CONNOR, D.G., Influence of HeatingRate on Transformations in an 18%Ni Maraging Steel,Nature, 213, 170 (1967).

6. Ni-Co-Mo MARAGING STEELS —THE CONVENTIONAL GRADES

[6.1] LEOER, M.T., and AYMARD, J.P., Contribution to the Metal-lurgical Investigation of a Cast Maraging-Type Steel, Rev.Met., 68, 783 (1971).

120

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OH M l --I ONI U M M , H l u l i S I R I \ ( i [ H S H I T S

"\ \i-(.t>-Mo MAUU.IM. Sll-.l.l.NiY. IXTR\-HU;H SlKIMiTH tiRADKS

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['.?) MM.NII . -V. \ m m R. P.. PKAI'IIK. J.M.. C"i ISOI HM>IS, 1) .and H\mi4Vis. 1 . Microslructure. Strength and Toughnessof l.'NiUOtli Maraging Sleel. Cobalt. m73. 3

P ' I ] SSM'I. 1 . and \ H I R V h.J.. The Properties of INNi (350)Maraeini: Steel ProJticcd fri.'m 1 Icmental and Prcallo^cdPowder*!" i;-«Je, Met. 15. 332 (W2i

[""I CXRIIR. C.S.. The I ifei.1 of He.it Treatment on ihe Fractureloughness and Subcntical Crack Growth Charscieiisiics ofa 3*l'i-Oradi- Maiaging Steel. Met. Trans.. 1. 1551 (19"UI.

[". j HINRY. R.J.. and C*R>. R.A.. An Fxtra-Strong MaragingSteel. Metal Pn'«r.. 96 (3i. 12* I!9h9>.

S«K, H.J . and K^Li.sH. D.. Improved Fatigue Resistanceof 18Ni 13501 Maraging Steel Through ThcrmomcchaniealTreatment*. Mel. Tram.. 5. <•«:> (iy"4|.

[•1.1]

[S.j]

8. STAINLESS MARAGING STEHLS •-PHYSICAL METAI.MRGY

BRH.C.S. J.Z.. and PARKLR. T.D. " The Super 12",,Cr Steels " .Climav MoKbdenum Co.. New York, lflds.

Pu\n\. J.. and V m i n . P.. Contribution to ihe Study ofTernar\ Steels Conunnine Chromium. Chiinie lihlmlri?,Numer'o Special. 404 (W;u>"

CorrsovRADis. D.. The F.tTect of Cobalt Additions to Precipi-tation-Hardening Steels. Man. Sii. Re\. Met.. 58. 503 (I9f>||.

i. F.B.. Con-Iron Steel hist..

[.V.Vj iRviNh. K.J.. LLF.WILI.VS. D.T.. andtrolled-Transformation Stainless Steels. /192. 218 (19591.

[#.5J HAMMOND, CM., The Development of Maraging StainlessSteels Containing Cobalt. Cobalt. No. 25. 195 (1964).

[8.6] BL'NGARDT. K.. SPYRA. W.. Contribution to the Developmentof Stainless Maraging Steels, DEW Tech. BIT.. 9. 361 (1969).

[S.7] ALLT. R.T., HOLTMANS. R.B.. and MYERS. J.R.. Heat Treat-ment of a Cr-Mo-Co Martensitie Stainless Steel for OptimumCombination of Strength. Toughness and Stress-CorrosionResistance. Trans. A.S.M., 61. 75 (1968).

[8.S] DLTIS, E.J.. and HABRAKES, L.. Precipitation-Hardening Stain-less Martensitic Steels. Paper presented at "" 41' Congres de laCommunaute Europeenne du Charbon et de 1'Acier ", Com-mission des Communautes Europeennes, 1968.

[«.P] WEBSTER, D., Development of a High-Strength Stainless Steelwith Improved Toughness and Ductility. Met. Trcts.. 7. 2097(1971).

[<S.IU] WEBSTER, D., Optimization of Strength and Toughness inTwo High-Strength Stainless Steels, Met. Trans.. 2, 1857 (1971).

[8.11] DIDERRICH, E., C'JUTSOURADIS, D., and HABRAKEN. L., VeryHigh Strength Stainless Steels, in " Journees Internationalesdes Applications du Cobalt, Brussels, 9-11 June 1964",C.N.R.M.-C.l.C, Brussels 1964, p. 318.

I*./-'I l.AMUiRt. N.. DiiM'ttu. J.M., ;ind CUITSIH IIADIS. O., Precipi-taimn-llardenini; I obalt-l. opper-MoKhdenum Stainless Steels,( .•halt. No. 47. M (1970).

(.""'./.'I (. AIUN. R I ., and M.WIAK. Ci N . Struclure and Properties ofa Neu High-Strength Maraging Stainless Steel, Cohalt, No. 55.y; (i-)^;i.

[S 14] l u m i M w . G.H . and Hi n . l-.C. ihe 1 Ifeci of Compositionon the 'temperature of Spontaneous Transformation of Auste-nue to Mariciisne in IS-M T'\pc Stainless Sleel. Trait.'.. .•t.S.M.,45. " (I453i

\J.I>] I I I I R I I S . S.. "Ihe Properties of I.o«-Carhon Iron-Nickel-Chronnuni Marlensitcs. Trum. I / A//.., 236. I4M (I9(>M.

[S. 16] 1K\ ISI . K.J.. I he I licit of Cobalt in Sleel. in •• Journees Inter-nationales des Applications du C oball. Brussels 9-11 JuneI9M ". C.N.R.M.-C.l.C.. Brussels I9M. p. 2M>.

IX.I"] Bi M.ARIII. K... Si'VRA. \V.. and 1.INSERT/. CI., lnvesiigalionsof Martensitic Nkkel-Chromium and Coball-ChroiniiiniSteels, in •• Siecl-Sirenpthening Mechanisms'", Climax Molvh-dentim. Zurich. 19^0, p. 4^,SMiowski. i P.. l)e\clopmcnt of a Stainless Marumng Siwl,\ksati En? Quart.. U (2), 47 (1972).WmsriR. D., Increasing the Toughness of the MarlensitieStainless Steel Al-C-77 by Control of Retained AusteniteContent. Ausforniing and Strain Ageing. Trans. A.S.M.. 61.

(.•v./.S'l

[X.IV]

[.S.20]

I.V.2/]

[X.2-]

\ii.24)

\X.25]

[X.26\

[S.27]

[8.28]

[..'29]

[$.30]

[SJ1]

[8.32]

[8.33]

WEBSTER, D.. The EITcct of Deformation Voids on AusteniteGrain Growth in Steels, Trans. ASM.. 62, 470 (1969).

FLLIOTT. R.P., " Conslitution of Binars Alloys". First Sup-plement. McGraw-Hill Book Co.. New >ork, 1965, p. 346.

WILLIAMS. R.O., anil PAXIDN, H.W.. The Nature of Ageing inBinary lion-Chromium Alloys around 500 C, J. Iron Sleelhist.. 185. 358 < 1957).WILLIAMS. R.O.. Further Studies of the Iron-ChromiumSystem, Trans. A.I.M.E., 212. 497 (1958).Pcivrv, G..and B\sni*.', P.. Transformations in Iron-ChromiumAlloyr. Near the Equiatomic Composition. Rev. Mel.. 53. 147(195h).ANTONY. K.C.. Anting Reactions in PrecipituUon-HardenableStainless Steels, J~Metals. 15. 922 (1963).

DULIS. E.J.. CHANDHOK, V.K., and HIRTH, J.P., Relationshipbetween Fatigue and Damping Characteristics and Microstruc-ture of 12%Cr Steels, Trans. A.S.M., 54. 456 (1964).

BLACKBURN, M.J., and NUTTING. J.. Metallography of anIron - 21 ".„ Chromium Allov Subjected to 475 C Embrittlement,J. Iron Steel Insl.. 202. 6)6 (196-4).L A G N E B O R G , R.. Metallography of the 475 C Embritt lement inan Iron - 3 0 % Chromium Alloy, Trans. A.S.M., 60, 67 (1967).

M A R C I N K O W S K I , M.J.. FISHER, R .M. , and SZIRMAE, A., Effectof 500 C Ageing on the Deformation Behaviour of an I ron-C h r o m i u m Alloy, Trans. A.f.M.E., 230 , 676 (1964).

KOSTER, W., and H O F M A N N , G. , Equilibrium Phenomena in theTernary System I rcn-Cobal l -Chromium, Arch. Eisenhiitlenn:,30. 249 (1959).

COUTSOURADIS, D. , DRAPIER. J .M. , DIDERRICH. E.. and H A B R A -KEN. L.. Precipitation Hardening of High-Strength StainlessSteels. Cobalt. No. 36, 144 (1967).

Cot—sctiRADis, D. , DRAPIER, J .M. , D I D E R R I C H , E., and H A B R A -KEN, L., Role of Cobalt in the Strengthening Mechanisms inthe AFC-77 Steel A F M L Tech. Rept766-243, 1966.

P I N A R D - L E G R Y , G., D A C L N H A BELO, M. , MONTUULLE, J.,and C H A U P R O N , G. , Contr ibut ion to the Development ofSpecial Steels with High Mechanical Strength and Corros ionResistance, Mem. Sci. Rev. Met., 68 , 81 (1971).

122

RLFF.UIiNCF.S

\x..U] DmiKRicn, 1: . DKAIMIR. J.M.. (OITSOI KADIS, 1)., and HAIIBA-KI N. I... Role of CHhall in the Strengthening Mechanisms in theAIC-77 Steel, ,:; ML Tech. Rept. 65-425, I96.S.

\Kjf] 1 HoMisos, FA., and Wist. O.R.F., Inlermetallic CompoundPrecipitation in an I•e-IO",,C'r-13",,t'o-5"uMo Alloy, J. IronSleet /».»/.. 210, 69| (19721.

(S.MJ Wisr. O.R.I-'.. The Constitution of Iron-Rich Alloys of theFe-Cr-Co-Mo System : A Review. Cobuli, No. 51. 77 (1971;.

IK..)"] Pi ARSON. W.B.. " A Handbook of Lattice Spacings and Struc-tures of Melals and Alloyv. ". Pergumon Piess, London.Vol. I, I95N. and Vol. 2. 1967.

[f.JS] KINOSHITA. T , TciKfNAiiA, Y., and TOVOSIIIMA, T.. Oil theStructural Diagram anu Age Hardening of Maraging Stain-less Steels. Sipptm Kin:ukii (iakkui-Si, 33. 254 (19691.

[8.Jv\ KASAK, A.. CHAMHIIIK. V.K.. and Di us. F.J.. Developmentof Precipitation-Hardening Cr-Mo-Co Stainless Steeis, Trans.ASM.. 56, 455 (1963).

IS.401 Dvsos. D.J.. and K.EOWN. J.R.. A Study of Precipitation in aI2"nCr-Co-Mo Steel, Ada Mel., 17. 1095 (1969).

[8.41] CASTAONE. J.L., Development of Maraging Steels wilh ImprovedProperties, in " Steel-Strengthening Mechanisms ". ClimaxMolybdenum Co., Zurich. 1970. p. 89.

[8.42] HORNBOUFN. E.. Clustering in an a Iron-Molybdenum SolidSolution. J. Appl. Phys.. 32, 135 (1961).

[8.43] K.OUTSKV. J.. and JHEK. J., Composition of Precipitates inModified 12"uCr Steels in the Range above 550 C, J. IronSteel In.it., 200, 9.18 (1962).

[8.44] McMi'LLiN, J.G., RIITER. S.F.. and tatLiNG. D.G.. Equili-brium Structures in Fe-Cr-Mo Alloys, Traits. A.S.M., 46.799 (1954).

[8.45] CATON, R.L.. A Maraging Stainless Steel fcr H00 to 1100 FService, Metal Progr., 92 (1), 106 (1967).

[8.46] WFBSTER, D.. The Improvement of Mechanical Properties of.\Ff "7 by Austenite Grain Refinement. Trans. A.S.M.. t>2,759 (1969).

9. STAINLESS MARAGING STEELS —PROCESSING AND PROPERTIES

[9J] " AFC-77 "", Alloy Digest SS-209 (May 1968), EngineeringAlloys Digest Inc., Upper Montclair, N.J.

[9.2] " PyrometX-12 ", Alloy Digest SS-152 (Feb. 1964). EngineeringAlloys Digest Inc., Upper Montclair, N.J.

[9.3] " Pyromet X-15 ", Alloy Digest SS-242 (Sept. 1970), Engineer-ing Alloys Digest Inc., Upper Montclair, N.J.; "CarpenterPyromet X-15", Carpenter Technology Co:p. TechnicalNote, May 1969.

[9.4] " Carpenter Pyromet X-23 ", Carpenter Technology Corp.Preliminary Technical Data Sheet, 1973.

[9.5] BiJNOARDT. K., and SPYRA, W., Contribution to the Develop-ment of Stainless Maraging Steels, DEtf-Teeh. Ber., 9, 361(1969).

[9.6] BRANDIS, H.. and VON DEN STEINEN. A., D.I'.W., Krefeld,Germ..;iy. Private communications, 1973 and 1974.

[9.7] DULIS, H.J., PF.RLMUTTER, I., and KASAK. A.. Heat-ResistanlFracture-Tough Stainless Steels, WESTEC Pii.rer No. W9-3.3presented at 1969 Western Metal and Tool Conference andExposition, Los Angeles," Calif., March 10-13, 1969; KASAK. A.,Private communication, 1973.

[1.8] WIHSTFR, 11.. Increasing the Toughness of the ManensiticStainless Steci AFC-77 by Control of Retained AusteniteContent. Ausformmg and Strain Ageing. Trans. A.S.M.. 61,816(1968).

[y.9l WIBSTIR. D.. Optimization of Strength and Toughness inTwo High-Strength Stainless Steels. Met. Trans.,2. 1857(1971).

[V.KI] WIIISTF.R, D.. Oe\elopment of a High Strength Stainless Steelwith Improved Toughness and Ductility. Met. Trans 2 1097(I97|

[V.ll] Bt.-NtiARDT. K.. SI"»RA. W.. and LKVSAHT/. G.. Investigationof Ma;tensitic Ni-Cr and Co-Cr Steels, in •' Steel Strengthening,Mechanisms "'. Climax Molvbdenum, 7urich. 1969. p. 108.

[9./JI CHANUIIOK. V.K.. -\n<S KASAK. A.. Stainless Steel Wire FeaturesGreat Strength at High Temperatures. Meial Progr.. 91 (li,108(1967).

[9.IS] Dius. EX. and HABRAKFN. L.. Precipitation-Hardening Stain-less Martensitic Steels. Paper presented at " 4 " Congres dela Communautc Europeenne du Charbon et de 1'Acier".Commission des Communauxes Europeennes. 1968.

[9.14] WRBSTFR, D. Mechanical Propetties of AFC-77 StainlessSteel Bolts. Metals Eng. Quart.. 13(2). 53 (1973).

19 15] COLTSOURADIS. D.. DRAPIER. J.M.. DIDLRRICH. E.. and HAHRA-KIN, L.. Precipitation Hardening in High-Streneth StainlessSteels, Cobalt. No. 36. 144 (1967).

[9.16] LAMBERT. N.. DRAPIER. J.M.. and COUTSOLRADIS. D.. Preci-piiation-Hardening Cobalt - Copper - Molybdenum StainlessSteels, Cobalt. No. 47, 68 (1970).

[9.17] CATON, R.L.. A Maraging Stainless Steel for 800 to 1 IOO'FService, Metal Progr.. 92 (1). 106 (1967); CATON. R.L.. Privatecommunication. 1973.

[9.18] HAMMOND. CM., The Development of Maraging StainiessSteels Containing Cobalt, Cobalt, No. 25. 195 (1964).

[lU9] HAMMOND. CM.. AM-367. A Stainless Steel that CombinesHigh Strength with Notch Toughness. Cobalt. No. 18. 8 (1963).

[9.20] CATON. R.L.. and MANIAR. G.N.. Structure arid Properties ofa New High-Strength Maraging Stainless Steel, Cobalt. No. 55.92(1972).

[9.21] AULT. R.T.. HOLTMANN, R.B.. and MYERS. J.R.. Heat Treat-ment of a Cr-Mo-Co Martensitic Stainless Steel for OptimumCombination of Strength, Toughness and Stress-CorrosionResistance, Trans. A.S.M., 61, 75 (1968).

[9.22] WEBSTER, D.. The Stress-Corrosion Resistance and Fatigui;Crack Growth Rate of a High-Strength Martensitic StainlessSteel AFC-77, Met. Trans., 1, 2919 (1970).

[9.23] CATON, R.G.. and CARTER, C.S.. Evaluation of AFC-77 Mar-tensitic Stainless Steel for Airframe Structural Applications,Air Force Materials Laboratory Tech. Rept. AFML-TR-73-182,Sept. 1973.

[9.24] Wu, K.C, and KRINKE. T.A.. Weldability Studies for AFC-77.Part 1. Mechanical Properties, Welding J., Res. Stipp!..47, 332i(1968).

[9.25] WEBSTER, D., Optimization of Material Properties to AchieveStructure Efficiency Metals Mater., 5, 257 (1971).

[9.26] SPAHK. H., and KOPP. L.. Possibilities of Application of VeryHigh Strength Maraging Steels in the Chemical Industry.Chemie-liig'iueur-Tedimk, 41, 1243 (1969).

10. CONCLUSIONS

[10.1] ROSENFIELD, A.R., and MCEVILY, A.J., Some Recent Develop-ments in Fatigue and Fracture. Paper presented at 37thMeeting of Structures and Materials Panel, AGARD, DenHaag, Netherlands. 7-12 Oct. 1973.

123

I U'MAINIM. HIGH-STRENGTH STEELS

AUTHOR INDKX

\HRMi-Msos . r . i v . 11. I

M.i I N . Ci.H.. ti« (ft./2|.

• \ s sn . i . . G.S. . P [_\5t>|.

A M ' i M . K.I".. 1'1 I-V.-M

M-Ki<iinvi. i > . : : u-'-\i i r . R .T . . 24 [

.W j j . / 4 ) . "1y4 f.i'. 'i. i l l W.

• \ 1 M A R H . J . P . >'N

\t>.->\. ' '9

4 3 [ 4 . 1 4 ) . 4 hS4 [.v.-j. 9K

BBAKER, A . J . . 19 [2.6_\ .\rt.?]. 20 (j.nJ:.

58 [5.39]. 59 [5.49], 60 [5.-/V]. h l [5.-/9I.

BASIM. W.R., 58 [5.*)]. 59 [5.401

BVNFRJIE. B.R.. 58 [5.-/5J. 59 (5.5/1 M [.'..'/.5.6/). 62 I5.-/5. 5.5/. 5.(5/].

B A K D M . F.J.. 6S [6."\.

BARSfiM. J.M.. 4" [4.21). ^4 [ft.j 'i.BARTON. C.J.. iO [2.30], >R [5.3(5], 60 [5.3(5].BASTII-N. P.. 91 [tt.24].

BAYF.R. A.M.. 69 (5.8). 77 [6.8).BECKER. P.C.. 49 [-OS].BELL. T.. 10 [2.321. 12 [2.32\. 13 [ : . i i ) .BHAT. G.K.. 19 [2.6«).BIF.BFR. C.G.. 50 [5.J. 5.-/]. 54 [5.201.BILBY. B.A.. 8 (2.201-

BlRKLt. A.J.. 22 [3.7]. 28 [3.7). 29 [3.7],31 [3.7], 37 [3.7], 40 R.4], 42 [4.4]. 69 [(5./0].

BLACKBURN, M.J., 91 [8.271BOESCH. W.J.. 54 [J./9]. 53 [5.19], 77 [7.21.

78 [7.2], 7 , '..'), 80 [7.2].

BONIZEWSKI, T.. 55 [5JO).

BONNET, M.. 74 [6.2H].

BOIJLGER, F.W.. 76 [6.3*].

BOURCEOT. J.. 55 [5.28], 62 [5.28). 63 [5.281,

64 [5.28).

BRADLEY, A.J.. 50 [5./].

BRANDIS. H.. 102 [9.6], 103 [9.6). 107 [9.6],108 [V.6). 109 [9.(5], 110 [9.6J, 111 [9.6],112 [9,(5j.

BRIDGES. P.J., 57 [5J5], 66 [5.35].

BKIGGS. .I.Z.. 82 [«./], 83 [*./]. 98 [S.I].

BRisBANt. A.W.. 71 [6.20].

BROOK, R.. 69 [6.12].

Bui, NAM, 65 \5.76).

BL'STGARDT, K.. 84 [8.6], 86 [8.6]. 87 [8.6],88 [8.5. 8.17], 89 [S./7], 9i [8.17], 93 [S./7],94 [«.<}, 8.17]. 95 [S.6], 97 [S.6, 8.17],102 [9.5], 105 [?.J, 9.11], 106 [9.5], 107 [9.5,

v / / ] 109 [^.5]. 110 [V.5. V./i'l. I l l- . / / ] . M3 I'-5].

Bi s i l l . C.C . "I 1<S./7. (S.l^l.Brsti. M.!-.. 5 1.7./3). h I-1./.)].

CAIKNS. R.l . 9 [.r.M]. 54 [5.IS). 56 [5./.V|,h'J [ft.//]. " [ ' . / ] . 7S[7./i . ^9 [-./]. .SO [7./].

CAMI 'BUL. J.L.. "S |»-""|.

CAI'I-NIIS. J .M.. M )5.ft/]. 62 [5.ft/|.CARTiR. i:.S.. 4S [4.23). "5 lrt.3J[. 79 [7.71.

SO ['.7|. 112 [9.23].C A R I . R.A.. SO [?.!<].

CASTAONF. J.L.. 52 [5. VI. 95 [?t.4!l'C\TOS. R . G . . 112 l'J.23).

CATON. R.L.. 87 [«./.?]. R8 [«./.!). S9 {«./3].95 [cV.73], 96 [A',/31. 97 [8.45]. 110 [V./7].Ill [9.20].

CHANWOK. V.K.. 17 [2.53]. \S [2.59]. 38 [3.22].39 [3.22. 3.23], 91 [o..'i5], 95 [«..?W|, 96 [«.39],97 [».3y]. 98 [.«..f9!, 99 [S.39]. 1.06 [V.i2'\.

CHAI/DRON. G.. 94 [«.33|, 95 [«..'.»]. 97 [S.J3),99 [,?.33].

CHENG. I-LIN. 20 [2.65], 52 [5.12].

CHILTON. J.M.. 10 12.30], 58 [5.36], 60 [5.Jtf].CHW-TIAS. J.W.. 7 [2./5], 8 [2.20).CIZEROS, G.. 8 [2..'2], 11 [2.22), 65 [5.77. 5.HO].CLAUSING. D.P.. 42 [4.12], 47 [4.12].COFFEY. F.J., 75 \fi I ' ;COHEN. J.B.. 64. [5.70].COHEN, M.. 8 [2.21]. 27 [.?./•/, 3./5], 28 [ i .«I ,

29 [3./5I. 30 [3.15], 43 [4.17], 44 [4-/7],45 [4:/ 7, 4.18], 54 [5.22].

CONRAD. H., 20 [2.66).CORN, D.L., 73 [6.24].COLTSOURADIS, D.. 2 [2.2, 2.3), 7 [2.3],

9 [2.27. 2.28], 15 [Z27], 16 [2.27], 23 [5.//],24 [3.//]. 25 [3.11], 2fi 13.11], 28 [3.//],29 [3.//], 30 [3.//], 49 [4.27], 51 15.7],58 [5.44. 5.46]. 59 [5.44). 60 [5.53], 61 [5.-/(5],62 [5.-A51. 64 [5.7], 65 [5.7], 77 [7.5], 78 [7.5],79 [7.5), 83 [8.3], 84 18.3), 85 [8.11], Sft [S./2],87 [S.3, S./.2], 92 [8.3/], 93 [A'J2], 94 [8.11,S.32, 8.34), 95 [8.31, 8.34], 96 [8.J/],97 [8.11, 8.12], 98 [,S.//], 99 [8.31, 8.32],110 [9.A5, 9./rt].

COWAN, T.W., 54 [5./9], 58 [5.19]. 11 [7.2],78 [7.2], 79 [7.2], 80 [7.2J.

Cox, T.B., 28 [3./i5], 29 [3./6J.CRIMMINS, P.P.. 53 [5./3].

CROOKER, J.W., 47 [4.22].

CROWE, D.J., 4 [2.8].

CUPP, C.R., 6i [5.60], 62 [5.60], 63 [5.60],64 [5.60].

DOAIIKIW-SKI. O.S.. 22 |3 . . \ 3.6, 3.9), 28 [3.9].

31 I3.A). 32 [.*.y]. 33 [3.«]. 34 [3.9], 35 [3.9],36 (3.V). 17 [3,rt. 3.9], 38 [3.9]. 39 [3.9!,42 [4.11). 69 [I5./W].

DAnosi. F.. 65 [5.7(5].DA Ci'NHA Bi:i.n. M.. 94 [«..?3]. 95 [8.33\.

97 [N.33], 99 [S.33J.DAS, S.K.. 12 [2.39], 15 [2.39\.DAITUVICH, D.P.. 74 |6.29, 6.3(.»], 75 [t>.34],

76 [6.29, 6.34).DAVIIS. R.G., 14 [2.^5J.DECKER. R.F.. 51 [5.6]. 54 [5.14. 5.IM

60 [5.6. 5.56], 62 [5.56], 65 [5.75], 73 ]6.2I).DETERT. K... 20 U.6^], 59 [5.50], 62 ,r5.i5.?].

65 [5.63], 66 [5.50. 5.63].DEVRIES, R.P., 71 [6.15],DIDERRICH. E., 60 [5.53], 85 [S.//]. 92 [8.31],

93 [8.3.'].94 [S.//, S.3.', 8.34], 95 [S.3/, S.3J],96 [8.31], 97 !».//]. 98 [*.//]. 99 [8.31, 8.32),110 [9./5].

DIENTS. G.J.. 15 [2 5/].DlRAN, L.M., 71 [6.14).DRAPIER. J.M.. 23 [3.//]. 24 [3.11], 25 [3.//1,

26 [3.//], 28 13.//]. 29 [3.//]. 30 [3,11).49 [4.27], 51 [5.7], 58 [5.44. 5.46], 59 [5.44],60 [5.53], 61 [5.46J, 62 [5.-W1, 64 [5.7],65 [5.7], 11 [7.5]. 78 [7.5]. 79 [7.5], 86-.S./2].87 [8.12], 92 [«.3/]. 93 [8.32]. 94 [S.32, 8.34],95 [5.3/. 8.34]. 96 [S.3/], 97 [8.12], 99 [8.3/,A'J-'l, 110 [9./5, 9.16).

DIJBOIS, B., 65 [5.78].

DULIS, E.J., 17 [2.53], 18 [2l59], 38 [3.221,39 [3.22. 3.23], 84 [5.8], 89 [8.8], 91 [8J6J95 [8.39], 9fi [S.39], 97 [«.«, 8.39], 98 [8.39],99 [8.8. 8.39], 102 [9.7], 103 [9.7], 104 [9.7],105 [9.7], 106 [9./3], 107 [9./3], 108 [9.7],109 [9.7], 111 [9.13], 112 [9.7, 9.13],113 [9./3J.

DUNNE, D.P., 7 [2./6J.

DYSON, D.J., 95 [8.40], 96 [8.40], 98 [8.40],99 [8.40].

EASH, J.T.. 51 [5.6], 60 [5.6].EBELING. D.G.. 96 [8.44).

ECONOMOPOULOS, M., 4 [2.9], 5 [2.9], 6 [2.9].EICHELMAN, G.H., 87 [8.14].

ELLIOTT, R.P., 91 [8.21].

ELSEA, A.R., 36 [3.20], 40 [4.1], 42 [4.1],

48 [4./], 49 [4./].ENTWJSLE, A.R., 7 [2./9J.

ERICSSON, T., 64 [5.70].

124

AUTHOR INDEX

I--H-LDIN-O, J., 77 [6.41].FINE, M.f.. 16 [2.55J, 17 13.55]. 58 [J..?7].

59 [5.37]. 60 [5.37]. 63 (.f.rtrt). 64 [5.ft9].FlSHEK, R.M.. 91 [8.29J.FI.INV. I1 A., 2 |2.4J.FiiiIU.IN. S.t !4 [2-46. 2.48. 2.49]. 15 [2.41-i.

3.M\. 20 [-'.Jrt]. 37 [.(.2/1, 52 [5.8]. 54 [5.7.5.5.2/|. 58 [5.8|, 60 |5 .2/ . 5.56]. hi [.?._'/. 5.5VJ.1.2 15.X. 5.56. 5.59]. 63 [5.ft. 5.65]. 65 [.57.fl.(•7 [5.8. 5.31. 5.5V]. 73 [ft.2/]. 74 [6.3V].1? [6.34]. 76 [6.29. 6 J 4 ) . HI [,S./5], 94 \».I5].1(11 [«./.?].

tiARVvocm. R.D. . 5K [5.43], 59 [5.4;!]. 60 I5.43].CiiiiiuiNS, D.F . 62 [5.62].

GILBERT. A.. 55 15..M).OILI-.WKV, K.P.. 70 [6.78], 71 |6./8].

GiuniiERG, A., 67 [J.AV].GOLDMAN. A.J.. 9 [.'.29], 51 [5.(5], 60 [5.6].

GOLDSCHMIDT, H J . , 50 [5./J.GRAAE. A.. 74 [6.271.ORANGE, R.A., 9 [2.2.?].GRAY. H.R.. 76 [6.36].GREEN, S.J., 35 [ i . / 9 ] .GRW.ST.VEUJ, T.P., 24 [3.I3\. 36 13.20],

40 [4.1], 41 [4.70]. 42 [4./]. 47 [4.10].4S [,<./], 49 [4.!].

GROSS, J.H.. 22 [5.4. 3.6]. 31 [3.6]. 37 [3.61.42 [4.13].

GliLBRANSEN, L.B.. ">S [6J51-

GLRKLIS. J.A.. 76 [6.38].

HHAURAKEN, I... 2 [2.2. 2.3]. 4 [2.9]. 5 [2.9].

6 [2.9. 2.I4\, 7 [2J ) . 9 [2.28]. 23 [3.7 7],24 [3.111 25 [ 3 . / / 1 26 [3.//J. 28 [3.11],29 [ J . / / ] , 30 [3.11], 49 [4.27], 51 [.5.71,60 [5.53], 64 [5.7], 65 [5.7] 77 [7.5], 78 [7.5].79 [7.5]. 84 [S.5], 85 [8.11], 89 [8.8], 92 [8.37],93 [SJ2], 94 [8.77, 8.52, S j f l . 95 [8.31.8.34],96 [8.37], 97 [S.S. 8.77], 98 [5 . / / ] , 99 [S.8,S.3I. 8.32], ' 06 [9.13], 107 [9.73], 110 [9./5],111 [9.13\. 112 [9.73], 113 [9.13].

H,\LL, A.M.. 1 [ / . / ] . 36 [3.20], 40 [4.7].42 [4.1], 48 [4./] , 49 [4.7], 50 [J.5], 77 [6.40].

HAMAKER, l.C, 69 {6.8], 77 [6.8].HAMMOND. C M . , 9 [2.26], 56 [5.32], 71 [6.75].

84 IS.J], 86 [8.5], 87 [8.5L 110 [9.78, 9.79].

HASAN, S.Z., 75 [6.35].HAUSER, J.J.. 58 [5.45], 59 [5.57]. 61 [5.57,

5.61], 62 [5.45, 5.51. 5.61].HAWN. J.M.. 71 (6.20}.

HAYNES. A.G., 77 [6.41].HAYNES, M., 74 [6.28].

HENRY. R.J., 80 [7.8].HENTHORNE. M.. 75 [6.33].HIRTII, J.P., 18 [2.59], 38 [3.22], 39 [3.22.3.23],

91 [8.26]HOIMANN, G., 92 [8.30].HOLTMANN, R.B., 84 [8.7], 98'[8.7], 99 [8.7],

111 [9.27],

HONEYCOMDE. R.W.K.. 4 [2.7].HORNBOUEN'. E., 18 [2.5*]. 64 [5.68], 96 [8.42].HOTZLER, R.K.. 22 [3.10]. 32 [3.10],HOUDREMONT, E , 39 [3.24].HUIZING, H., 55 [5.25].HULIT, G.W., 68 [i5.2J.Hiî-L. F.C., 84 (footnotel. 87 [H.I4].

I

IAVNELLI, A.A., 77 [7.J1. 79 [7.3].IMHOF. E.J.., 47 [4.21]. 74 [6.25].

IRVINE. K.J., 2 [2./J, 3 [2.6], 4 [2.8. 2.10],5 [2.12]. 6 [2.70. 2.12]. 7 [2.72], 15 [2.7, 2.52].16 [2.7. 2.52), 19 [ 2 / 1 . 84 [8.4], 88 [8.16].94 [8.16]. 97 [8.16]. 98 [8.16].

JACK, K.H.. 18 [2.57].JACKSON. J.K., 62 [5.62].

JANA. S.. 66 [5.82].

IECEK. J.. 96 [8.4i].JONts. F.W., 50 [5.2].JONES, R.D.. 58 [5.4J], 59 [5.43]. 60 [5.-O].

KALISH. D.. 27 [3.14.3.15]. 28 [3.14], 29 [5.7J],30 [.?.7J], 43 [4.77]. 44 [4.17]. 45 [4.17. 4.18].60 [5.57, 5.58], 77 [7.4]. 78 [7.4], 79 [7.4].80 [7.9].

KARDONSKII. V.M.. 15 [2.50].KASAK, A.. 17 [2.J.?]. 95 [8.39]. 96 [SJ9].

97 [8.39]. 98 [8J9] , 99 [8.39]. 102 [9.7].103 [9.7]. 104 [9.7], 105 [9.7], 106 [9.72],108 [9.7]. 109 [9.7], 112 [9.7].

KAUFMAN, L., 8 [2.21]. 54 [5.22].

KELLY, A.. 16 [234], 17 [2.54].

KELLY, P.M., 5 [2.13]. 6 [2.13], II [2.34],12 [2.34. 2.58].

KENYON, N., 76 [6.39].K E O W N , J.R.. 95 [8.40], 96 [8.40], 98 [8.40],

99 [8.40].KESSLER, H., 66 [5.87].KINOSHITA, T., 64 [5.77, 5.72], 95 [8.38],

96 [8.38].KLOSTERMANN, J.A., 55 [J.25J.KONKOL, P.J., 22 [3.4, 3.5].

K O P P , L., 113 [9.26].K O P P I , W.A., 68 [6.5], 69 [6.5], 75 [6.3].KOSTER, W., 92 [8.30].KOUTSKY, J.. 96 [8.43].KOVESI. P., 19 [2.61].KRAUSS. G.. 7 [2.77], 10 [2.17], !1 [2.77],

12 [2.17], 66 [5.S5].KRINKE, T.A., 113 [9.24].K U L A , E.B.. 20 [2.67].K U L I N . S.A., 27 [5.74. i . / J ] , 28 [5.74],

29 [3. /J] , 30 [5.75], 43 [4.77], 44 [4.77],45 [4.77, 4.78].

KURDJUMOV, V.G., 14 [2.47], 15 [230].

LACOMBE, P., 8 [2.22], 11 [2.22], 65 [5.80].LAGNEBORG. R.. 91 [8.28].

LAMBERT. N.. 23 [5. / / ] . 24 [3.11]. 25 [5.77]26 [5.77], 28 [3.11]. 29 [5 . / / ] . 30 [3.11],49 [4.27], 58 [5.46], 61 [5.46]. 62 [5.46]86 [8.72], 87 [8.72], 97 [8.12]. 110 [9.16].

LANG. F.H., 76 [6..?s/J.

LANGE, E.A.. 47 ]4.22\. 71 [(5./6J.

LEAVERLAND. J.V., 19 [2.61].LEGENDRE. P., 52 [5.11]. 65 [5.79] 66 K 79]

76 [6.57].

LEGER, M.T.. 68 [6.1].

LENNARTZ. G.. 88 [8.77]. 89 [8./71 91 [8 17]93 [8.17]. 94 [8.17]. 97 [8.77], 105 [9.11].107 [9.7/]. 110 [9.1 i]. I l l [9.77],

LEOPOLD. C.T.. 24 [J./3], J! [4.70], 47 [4.10].

LESLIE. W.C.. 3 [2.5]. 13 [2..?i5], 14 [2.44].16 [2.56], 18 [2J6J, 35 [3.19].

LTTTLE, R.T.. 24 [3.73], 41 [4.10]. 47 [4.70].LLEWELLYN, D.T.. 3 [2.6], 15 [232]. 16 [2 51]

84 [8.4].

LOTJERING. G.. 18 [2.58].

LUTZ, J.L.. 58 [5.40]. 59 [5.40],

M

MCEVILY, A.J,. 114 [70.7].

M A C H U N . E.S.. 13 [2.42J.MACIAG. R.J.. 22 [3.70]. 32 [3.70].M C M U L U N . J.G.. 96 [8.44].MAGEE. C.L.. 14 [2.45].MAGNEH. A.. 51 [5.7], 58 [5.44], 59 [5.44].

64 [5.7], 65 [5.7]. 77 [7.5]. 78 [7.5].79- [73].

MAITREPIERRE. P.. 55 [5.28). 62 [5.28].

63 [5.28]. 64 [5.28].

MANENC, J., 55 [5.28]. 56 [5.33]. 62 [5.28],63 [5.28], 64 [5.28].

MANQANELLO, S.J.. 22 [3.6], 31 [3.6], 37 [3.6].MANIAR, G.N. , 87 [8.75], 88 [8.75], 89 [8.13],

95 [8.73], 96 18.73], 111 [9.20].MARCINKOWSI, M.J., 91 [8.29].MARCUS, H., 58 [5.37], 59 [5.37]. 60 [5.37],

63 [5.66], 64 [5.69].MARDER, A.R., 7 [2.77], 10 [2.77], H [2.77],

12 [2.17].MATAS, S.J., 22 [3.7, 3.2], 35 [3.2]. 36 [3.2].

40 [4.2, 4.3], 43 [4.3, 4.16], 45 [4.76],46 [4.76], 49 [4.28, 4.31]. 69 [6.9].

MAYNOR, H.W., Jr.. 71 [6.17, 6.19].MELNICK, L.M.. 58 [5.40], 59 [5.40].MIHALISIN, J.R.. 54 [5.20], 55 [5.23], 60 [5.55],MILLER, G.P. . 55 [5.29], 57 [5.35]. 58 [5.29],

59 [5.52], 60 [532], 61 [5.52]. 66 [5.35].MINER, R.E. , 62 [5.62].M..CHELL. W.I., 55 [5.29], 58 [5.29], 59 [532],

60 [5.52], 61 [532].MONTUELLE, J . 94 [S.33]. 95 [8.33], 97 [8.33],

99 [8.33].MOON, D.P., 68 [6.7].MOURIKIS, S.. 64 [5.70].

125

l-OBALT-CONTAlNINli HK.H STRFNGTH STFFl S

\ l ! V . I K . 1 1 1 ' . :•%,-:

M M R V J R . S 4 ( S '

111 i« : / i

'•• H M

\ l , IMIM.N. R II . Ir. 12 54|. I" |2.54|

\ . . \ \K. (. J . h'l I'. .'M. "1 I" 'J)N i r n v . . J . n I-'.'"•*I. l-"1 i - ' - ' J . - ' • ' ' ! •

21 | 2<w | . " I I* . " 1

O

O V O N M I R . O Ci . h" (5 ."«4|.

O K A . M . . u i : <<]

OKAMIT. . . M . ^S 15.J/]

O K I H . I".. 63 |5 64)

OLOISON, C" T . "6 |ft J,v|.

OLSON. J H . ^4 [ft ."'i|.

O W I N . W.S.. 10 J2 OJ. 12 |2. '21. 13 [2.32].?? [5.24. 5.27 |.

PARKI-R. T.D.. X2 [S 11 xi [H.l\. 9H [,f./|.PARR. J.G.. 55 15.26].

PwintR. J.S.. " [j./.v|. 12 \:.4i\. 22 [j.;i..'1 [J./7], 35 [J.2]. 3<- |J.2|. 43 [4.15. 4.M].45 |4./5. ././ft]. 4fi [4 /ft).

PASOLIM. O.L.. "4 [6.2A|.

P\rriRsov. R.L . 10 [2. J»]. 11 l-'.JJ]. «> !5>*'i.PATTHISON. VV.R.. 54 [5.17]. ?> [5./T|.PAI LINA. J P.. 69 [6.W|.

PAXTOS. H.W.. 69 [6.l.i\. 74 (A./.»]. 7? [6..)/),91 [#..V].

PAVSDN. P.. 9 [:.23\.

PEARSON-. W.B.. 5>< 15.47]. 94 |«..^7).PEISTRLC. J.N.. M (5/>6).

PERKAS. M.D.. 15 [2.50). 52 [5.V. 5.WJ.

53 [5.4*J, 58 (5.V]. 59 [5.10]. 67 15.9).PERLMI;TTFR. I.. 102 [9.7]. 103 [9.7). 104 [9.7],

105 [9.7], 108 [9.7], 109 [9.7). 112 [9.7],PERRY. T.E.. 22 [3.1]. 40 (4.2).PETERS. D.T.. 54 [5.16). 61 [5.60], 62 [5.60],

63 [5.60. 5.<5J. 5.157], 64 [5.60], 65 [5.74],66 [5.16. 5.741 67 [5.74], 72 [6.22].

PETTY. E.R.. 14 [2.43].

PICKERING. F.B.. 3 [2.6]. 4 [2.S. _'./0, 2.//],6 {2.10. 2.11]. 15 [2.52], 16 [2.52], 84 [S.4].

PIN'ARD-LEGRY, G.. 5X [5.42). 60 [5.42].66 [5.42]. 94 [«..?.»]. 95 (»..«]. 97 [SJ.?],99 [#JJJ.

PITLER, R.K., 58 [5.4«]. 60 [5.48).PITSCH, W.. 66 [5.«/J.

POMEY, G.. 91 [8.24].

POMEY. J., 83 [8.2).

POOLE. S.W., 22 [3.1. 3.3]. 40 [4.2]. 48 [4.26].PORTER. L.F.. 15 [2.51]. 22 [3.5. 3.6. 3.7. 3.9).

28 [3.7. 3.9), 29 [3.7]. 31 [J.rt. -?.7], 32 [3.9],33 [J.9], 34 [i.9], 35 [3.9], 36 [J.9], 37 [3.6,3.7. 3.9), 38 [J.9], 39 [J.9], 40 [4.4], 42 [4.4,4.11], 69 [6,70].

PRIESTER, L., 64 [5.73].

PK.K n « , R.P.M.. 69 [6.131 74 [6./.?].PlMi-HRfY. W.!.. 50 [5.2]Pi ZAK. PP., 71 [A.M|.

R

RAR('<«-K. SG . 35 {J./vl.

RAI-K. H.J.. t.0 [f.57. 5.fX\. 77 [7.4). 78 p.4),"9 ['.4), HO (7.9|.

RADCLIFFF. S.V.. 7 [2.IS). 12 [24 /1 .

RM.HAVAN, M . 12 [240].

RATHBUNT. A.M.. 22 [3.4. 3.5). 40 [4.5].

42 [4.51, 4S [4.5]. 49 [4.30].

REISDORF. B.C. 19 [2.62). 20 12.62]. 58 [5..«].59 [5.3$. 5.49], 60 (5.49), M [5.49).

RUTER. S.F.. 96 [8.44).

RICHARDSON. L.S.. 54 [5/7], 55 [5/7],Kits, G.D., 22 [3.3], 48 [4.26].ROBERTSON, VV.D . 9 [2.29J.

ROLFE, S.T.. 47 [4.21], 74 [6.25].ROMINT, HE., 48 [4.24).

ROSEN-FIELD. A.R.. 114 [10.1].

ROSEN-STEIN. A.H.. 28 [3.16\. 29 [ ?./«].ROTH. M.. IS [2.58].

SAOOWSKI. E.P.. 54 [5.14], 68 [6.3], 69 [6.JJ,75 [fi.il. 89 [8.18].

SAVAS. J.. 49 [4-25].

ScHMATZ. D.J.. 57 [5.34).SCHRADER, H.. 39 [3.24).

SCHWARTZ, L.H., 58 15.37). 59 [J..?r|, 60 [5.37].64 [5.69].

SERVANT, C , 65 [5.77],

SHAPIRO, G . 66 [5.83].

SMMIZU, K., 12 [2.37], 58 [5.41).SHYNE, ).C, 57 [5J4],

SIMONS, R.C., 24 [i./J], 41 [4./0], 47 [4.10].SKELTON, M.W., 75 [6.35\.

SLUNDER, C.J.. 77 [6.40].

SMITH, H.L., 22 [i.«], 29 [3.8], 41 [4.9],42 [4.9], 48 [4.24], 49 [4.9].

SNAPE, E., 78 [7.6].

SNTTSAR', V.r., 52 [5.10], 59 [5.70].SOBER, R.J., 35 [3.19].

SONNINO, C.B., 75 [6.J5].

SPAHN, H., 113 [9.26].

SPEICH, G.R., 10 [2.30, 2.31]. 13 [2.31. 2.36),16 [2.56], 18 [2.36], 22 [3.9], 28 [3.9], 32 [J-.91.33 [3.9], 34 [3.9], 35 [J.9], 36 [J.9], 37 [J.9],38 [J.9], 39 [J.9], 42 [4.11], 61 [5.59],62 [5.59], 67 [5.59).

SPOONER, S., 60 [5.57].

SPYRA, W., 84 [8.6], 86 [8.6], 87 [8.6], 88 [8.6,8.17], 89 [8.17], 91 [8.17], 93 [8.17], 94 [8.(5,5.77], 95 [8.6], 97 [5.6, S./7], 102 [9.5],105 [9.5, 9.7/], 106 [9.5], 107 [9.5, 9.7/j,109 [9.5], 110 [9.5, ? . / / ] , I l l [9.5, 9.11),113 [9.5].

STACEY, A.G. , 14 [2.43].

STAVROS, A J . , 75 [6.31].

ii.A.. 46 [4.19], 47 [4/9).73 [6.23]

STI.NISIFHI. F.R.. 22 [3.8], 29 [3.8], 41 [4.9).42 [4.9], 4<< [4 24], 49 [4.9]

Swiss . P.R.. 10 [2.31]. 13 [2 31], 19 [2.6J],*" [5.J9].

SWASSON. W.l) . 55 [5.26].

S/IRMAI. A.. 91 [,S.29i.

THIVEMN. J.-P.. 8 [2.22], H U'-'-'I. 65 [5.80)THIVHUKR, P., 5ft 15..?.*]

THOMAS. B. 55 [5.28], 56 [5.J?|. 62 [5.2*1.63 \f.28). 64 [5.2.1).

THOMAS, G . 12 [2.J9. 2.40], 15 [2.39], 20 [2.65],52 15./2].

THOMPSON-, F.A., 95 [8.35], 96 [8.35]. 97 [S.J5].TOKI-SAUA. Y., 63 (5.64], 64 [5.7/, 5.72),

95 |.V..M], 96 [iV.J«].TOYOSHIMA. T., 64 [5.71], 95 [8.38], 96 [#.J,S].TROIANO. A.R., 76 [6.J61.

TRUSCULESCO, M.. 33 [3.18].

TUFFN-ELL, G.W.. 9 [2.25). 54 [5.if.). 56 [5./*].74 [6.26], 77 [7./], 78 [7.1], 79 (?./). 80 [7./).

uU< HIDA. J.M.. 49 [4.29).

VELTRY, F.J.. 78 [7.6).

VIALATTE, B.. 65 [5.78].

VIATOUR. P., 51 [5.7], 58 [5.44,5.46), 59 [5.44],61 [5.46], 62 [5.46], 64 [5.71, 65 [5.7],77 [7.5], 78 [7.5], 79 [7.5].

VISHNEVSKY, C , 46 [4./9J, 47 [4.19]. 73 [6.2JJ.vox DEN STEINFN, A., 102 [9.6]. 103 [9.6].

107 [9.6], 108 [9.6], 109 [9.6], 110 [9.6],111 [9.6], 112 [9.6J.

VOULET, P., 83 [8.2].

wWARD, D.M., 57 [5.J51. 66 [5.35].WAYMAN, CM., 7 [2.76], 10 [2.JJ], 11 [2.JJ,

2.J5). 12 [2.37]. 60 [5.54]. 66 [S.82].

WEBSTER, D., 84 [8.10], 85 [8.9,8.10], 89 [8.79],90 [8.19, 8.20], 91 [«.9], 100 [8.9, 8.10, 8.19],101 [8.70,8.46], 104 [9.8,9.9.9.10). 105 [9./0],106 [9.8], 108 [9.8], 110 [9.74], 111 [9.22],112 [9.22], 113 [9.74, 9J.5].

WEST, D.R.F.. 95 [8.35. 8.36], 96 [8.35],97 [S.J5].

WILLIAMS, R.O., 91 [8.22, 8.23].WILSON, E.A.. 10 [2.32], 12 [2.J2], 13 [2.321

55 [5.27].WOSKOWIAK, A.J., 22 [3.10], 32 [J./O].Wu, K.C., 113 [9.24].

YF.O, R.B.G., 9 [2.24], 56 [5.J/], 57 [5.J7].

ZACKAY, V.F., 57 [5.34].

S26

SUBJECT INDEX

SUBJECT INDEX

lfC-77 Haimtic transformations iSre al\o Lowercold working, 113 bainite. Massive hainile and Upper bain-composition, 82. i««-"«- 4-7- 23-2K.heat treatment, 84-85 89-90. 98, 99-101. B. and B, temperatures. 23. 24

102, 103, 105. 110, 111-112.hot working, 102. 112. ( '«• -V«' AFC-260machining, 113. Carbidesmechanical properties, 98, 99-101, 103-110. formation. If-. 17. 1X-I9. 25-2". 29-33. 39.oxidation resistance. 112, 113. 9?. ^-9H. 99stress-corrosion resistance. 111-112. morphology. 2ft. 31. 32. 33. 3f. 44. 9Kstructure, 84,85,89-91,95,96,97,98,99-101. orientation. 5. 1ft. 25-26. 2V,. 30thermal stability, 108-109. C'arbide-sircngihened steels, general. 22-49.thermomechanical treatniem. 90, 101, 106, bainitic transformation. 23-2S.

109, 112. carbide formation. 29-33.uses, 108, 109-110, 113. martensilic transformation. 23-24. 2S-29.welding, 113. strength toughness r.v. siructure relation-

iFZ-//\ modified (Alloys A and B) ship. 38-39.composition, 82. Carbide-strengthened steels. HP9-4-X gradesheat treatment, 90-91, 100, 103. ' See also HP 9-4-20. HP 9-1-25. HP 9-4-30hot working, 112. and HP 9-4-4?)mechanical properties, 100, 101, 104-105. primary processing. 40-42.stress-corrosion resistance, 112. secondary processing. 48-49.structure. 90-91, 100, 101. uses. 49.

AFC-26Q (C5Q) Carbide-strengthened steels. Co-modified Iow-composiiion. 82, 100. alloy. 18. 19.heat treatment, 100, 102, 103. 105. Carbide-strengthened steels. Co-modified 5Ni-mechanical properties, 100,103-105,107, 108. Cr-Mostructure, 84, 85, 100. heal ireatmem. 29.

Alloying elements mechanical properties. 29. 31-32, 37.aluminium, !5, 52, 56. 84. 88. structure. 31.beryllium, 52. Carbide-strengthened steels. 10Ni-8Co-Cr-Mocarbon. 12, 13, 14, 24, 29. 32, 33, 34, 35. 38. composition. 22. 32. 35. 3ft.

47, 69, 74, 83, 84. 86, 97. heal treatment, 28. 29. 32-?3. 34. 35. 37-38.chromium, 12. 14, 23-24, 34. 36, 38, 39, 52, 41, 42. 43.

53, 56, 81, 84, 86-87. 91-98, 110. mechanical properties. 28, 29. 32, 34, 35.cobalt, 3-4, 7, 9, 12, 14, 15, 16. 18, 19, 23, 37-38, 42, 43. 45, 47. 48.

29, 31, 33, 34, 37-39, 51, 52, 53, 57, 59, melting. 40.60-64, 65, 67, 81, 83, 8*4, 86-87, 88, 92-99, stress-corrosion resistance, 47-48.110, H4. structure, 28. 32-34.

copper, 15, 52, 84, 87. 97, 110. uses, 49.magnesium, 69. welding, 49.manganese, 12, 36, 52, 54, 55, 69, 84, 88. Cold working, 48. 76, 78. 113.molybdenum, 15, 18, 23-24, 34, 36, 38, 39, effect on properties and structure, 71-72. 90.

51, 52, 53. 56, 59, 60-64, 65, 67. 84, 86-87, '05, 106.88, 94-97, 99, 110. Continuous cooling transformation curves,

nickel, 9, 12, 13, 14, 35, 38, 39, 50, 52. 54, 23-24.55, 56, 59, 62-63, 65, 67, 84, 86-87, 88, Controlled-transformation stainless steels, Co-92-94, 100, 110. modified. 15-16.

niobium, 52, 56, 90, 101. Corrosion resistance, 110-111.nitrogen, 18, 84, 86. Creep-rupture strength. 46. 47. 72. 73. 108.silicon, 19, 36, 52, 56, 69, 84, 88.tantalum, 52. D.70, 82.titanium, 52, 53, 56,65,67,75, 84,86,87,97.tungsten, 52, 84, 97. Elastic properties, 69, 78. 79, 103.vanadium. 18. 23-24, 36, 52, 56. 84, 88.zirconium, 52. Fatigue strength, 46. 47, 72. 74. 80. 109-110.

AM-367, 82. Ferrite (delta)Applications. See Uses. effect on properties, 100-101.Austenite reversion, 8, 32-33. 34, 50, 54-55, occurrence, 15, 83-85.

64-67, 88-89. Fracture toughness. See wider Toughness.1, and At temperatures, 8, 9, 24, 56-57, 88.

effect on properties, 66, 79-80, 94, 107. Grain size, 62, 85, 90-91.Austenite stabilization. See Retained austenite. effect on properties, 21, 28, 75, 90-91, 101.

Heat treatment. 40-42. 68. 77-78. 84-85. 102-103. 113.

effect on properties and structure. 14-15, 16,17, 18. 27-34. 35, 36. 37-38. 42-43. 71,74-75, 79, 80, 84-85, 89-91. 98, 99-101105. 109-112.

High-temperature properties. 4ft. 72. 73. 79-80,107-109.

Hot working, 40. ftK. 77, 102, 112.HP 9-4-20

composition. 22.heat treatment. 30, 41, 42-43.mechanical properties. 26. 27. 30. 42-43,

45-48.stress-corrosion resistance, 47-48.structure. 23-24. 25. 28. 29.uses, 49.welding, 48-49.

HP 9-4-25composition, 22, 36.heat treatment, 40-41.hot working, 40.mechanical properties, 35, 42-45.structure. 24, 29.thermomechanical treatment, 43-44.welding, 48-49.

HP 9-4-30composition. 22.heat treatment. 30-31. 41, 42-43.mechanical properties. 26. 27, 30, 35. 42-43,

45, 48.stress-corrosion resistance, 47-48.structure. 23-26. 29.uses, 49.

HP 9-4-45composition, 22. 3b.heat treatment, 27, 36, 40-41.hot working, 40.mechanical properties, 27-28, 35. 36, 42-48.melting, 42, 43.stress-corrosion resistance, 47-48.structure, 24, 27, 30.thermomechanical treatment, 43^5.uses, 49.welding, 49.

Impurities, 69, 71, 74.IN-763. 53, 54, 67.Intermetailic compounds, general

formation, 16, 17, 19-20, 57-60, 62-64, 91-97.morphology, 58.orientation, 58-59, 60, 63.

Intermetailic compounds, typeso-FeCr, 91-99.(FeCrbMo. 95.X(Fe-Cr-Mo) °1-97, 99.o-FeMo, 58-uU.Fe->Mo, 58-60, 94-97, 99.U.-FB7MO6, 58-59, 94.Fe2(Mo,W), 96.o-FeTi, 58-59.FeVTi, 58.Ni->FeMo, 59.

127

* HIGH STRENGTH STF.FLS

N : . \ l , . I • " >

N . . I . s v . . M - . . - 1 - . • ) '

R . l e - l i - M , > . ••»,.. • «R . M . . - 1 . ' - I • '. ' 1 4 . ' l " ' ! »

| s , i i i i e : " - a ! f a n s ! , - ! i i u t i o n t i m e s . 2 4 . s 4 . s s

. ' l e i ! on precipi tat ion i c u t ions. I I . Ml .

'V!|'il,'!,V\ ll'-l 1 . 2V ; - l - ; <

;T*HVI iii>. I *-i V ^~>'*'.-:cmpci.iiur;: pn ' tvu^v . 2<v I". 4(.-4". 2.

"•, ~n. i ir . | iHi.•»•.•: tuinneViolation. 4. \ 2 < -> . 2~

M.^'mnni.1. 4.s. ~<\ "v I I 'Maraging steels t \ i - ( o-Mei. genetal. -U-M

preci" ut ion reactions. 14-20.pnnuir> priVcwiNt:. fiS. ~"- -s

he.11 iriMtnient.' 4

1» strucuire.'-(I. I I -

lien ircjuncni. hX. "?nKvhanical pr(>pcrtics.mclnng. ''*^!^c>^-vlK•t•nst"^ rc'^^lxin

Maraiiing >icc|v. ISNuM criulc

hc.it treatment. <»rv "4-"^.niLvhaniLal prttpcruev 51. hH-f^. ^!-""mcliinu. "Ivlros-cittTcsmn rcst^lancc. 7 S ^ 6 .

Murasinu '•led-. ISNiO?0i uruilctumpoMiinn. 5.1.heat trcaimcnt. hi*. " I . "A-"15

mechanic;!I propcTtiev 51, f>N-"4.niL-luni;. " I .^!rcss-ci>rro-Hin resistance. "t?-7d.ÎrucIurc. 5 \ 5s. 5 (). ft;. ftft. 1,".iherrmniK-'dunical treutmen:. ~0-~2.

Muraginu steels. IKNifiOOi [;radccompiiMlion. ' 3 . ft1*. "4.heal treatment. h.S. " I . 74-"1?.mechanical properties. 51. hN-^4.meltinu. 7 1 . 74.pnvv-der n tc ta l lu rg) . ^X. 69.sln;ss-C(irro>!iin resistance. 74-7h.structure. 5K. fS-i.lhermomechanicai treatmen;. 70-7;.

Maraging steels. IKNil.isni gradecold utirking. 7 s .composition. 53. 54.heat treatment. 77. 79. SO.hot working. 77.machining. 7x.mechanical properties. 51. 7K-XO.melting. 77.powder metallurgy. 7X-7lJ.stress-corrosion resistance, HO.structure. M).thermomcchanieul treatment, 79, 80.uses. 80.welding. 78.

structure, s'l. t.i-1,4. <•*•uses. SO. II s

MLiraeme steels. SNI(S(XM i:radei.om[>oûu»n. ^ \ ^4mechanical propeinev. ^ I

Miiragm • steels. I^Si- 'K'o gradeômpitsilion. .*. Mniechanical properties. "2. " 'uses. ~2

M.11,11:1115; steels. I s \ i - | s i o ur.uleconipos'lion. s; . S4mech.iinc.il ptoperties. "2uses. " ;

Maramni; steels. | s \ i - l « t ' i> i;rade Siv

Marayini: st<-'els. l^NiOMn t u d c . \ 1 , MM.11 lensitK transformations .Set1 i;/\;> I aih

m.irtensfte and Iwinned martensiiei. "-lf>.- ' - - 4 . -S-2''. s(|. <4-s". SS-S".

W.. \l . and \ / i lempenitures. s. 4. 12, 13._"'. I-',. 5h-5~, d4. X4. Sh-S".

\lassive baimteformation. 4-hmorpholou\. 4. sproperties, h."", 2>*. 2~.

Melting. 4(i. i.S. ~~. 101 -102.elfect on properties. 42. 43. 4". " I . "4. ~ \

Ordering, hfi. <MOxidation resistance. 112. i i 3 .

Powder meiallurg>. M. ^S-"1).Precipitation reactions See Carbides and

Intermeialhc compounds.Pynmiei X-12

composition. K2.heat treatrrenl. 102.hot working. 102.mechanical properties. 103.

Pyromet X-15composition. 82.corrosion resistance. 110-111.heal ireatmem. 102, 105.hot working. 102.machining. 113.mechanical properties. 103-105. 107. 108.structure. 97.welding, 113.

Pynmier X-23composition. 82.corrosion resistance. 111.heat treatment. 102. 103. 105.hot working. 102.machining. 113.mechanical properties. 88. 103-105, 107.structure. 8". K8. 95. 97. 99.

Retained ausienite, 13. 23-24 27, 32-33. 34,65-fi6, 87. 88, 89-90, 105-lOfi.

effect on properties, fifi. 88, 89, 99-100,103-104. 105. 111-112.

Self-tempering, 28, 29, 39, 41.Stainless maraging steels, general (See also

AFC-77. AFC-77, modified (Alloys A andB). AFC-260, AM-367, D.70, Pyromet

\-;.' . P\n:nfi \ - l \ Psiomt-I \-:j, VltralIi'ii4i)l. I. In,i/,'i I -1(12. and ( Irmfirl 40.1 , |KI -1 13

austenile resersion. NX-S9ferine (ilellai formanon. X't-X*. 100-101.grain size. 90-91. 101m.iilensilic transformation, .S5-X".precipitation reactions. 91-99.pnm:'r> processing. 101-103retained auslemte. K9-90secondary processing. 113strength toughness n structure. 99-101.uses. 1 I 3

Strenglhenuig mechanisms. 2-21.hummc phase lraiisliniiiaiions. 4-". 2"-28.martenstlic phase transf Hinaltons. ' ' - I ' I . 38,

plastic straining. 20-21precipitation. H>-20. 3s. I , I . ,,-. 94solnl solution. 2-4, IV Hi. 3". 3X. (^). ill, fi7.

Stress-corrosimi resistance. 4"-4X. '>^. 74-76,K0. 111-112.

lenipenni: resistance. 15, In. 19. 29. 30-31.38-39."

Tensile propertiesductilits. 14. 2h. 2"1. 30. 42. 4f>. dft. d8-(S9.

71. 72. 7S-79. 103. Kl.i. |0h.strength. 7 . 13-15, 19, > . 2". 28-29. 30, 31.

32. 34. 35. 37.39. 42-45. 48. 51, 53. (.1, 66.hi, f>8-72. 7K-79. 8S. 97, 98-101. ,'03-10(i.114.

Thermomechanical treatment. 20-21.effect on properties and structure. 43-45,

70-72.79.80.90,97. 101, 105, 10o. IU9. 112! ouiihnrss

fracture. 15, 27. 28. 42. 44-47. 51, 67, 68-72.78-79. 89. 98-101. iO l -W. 114.

impact. 7. 15. 19. 26-3i. 35. ;i7-39, 42. 43,fi7, 68-IS9, 78-79. 99. 103.

notch tensile strength yield (ultimate tensilelstrength. 35, 3f\ 43. hX-h9, 7R-79, 98, 99,103.

Twinned martensiteformation, I I . 12-13, 55. 85.morphology, II.properties. 13-15.

Clirufon 401composition. 82.corrosion resistance. 110.heal treatment. 102. 109. 113.machining, 113.mechanical properties. 97, 103-109.structure. KK, 95. 97.thermomechanical treatment. 105-106.uses. 113.

Ultra),in 402. 82.Vlirafiirr 403

composition, 82.corrosion resistance. 110.heat treatment. 102.mechanical properties, 103-104, 107, 109.stress-corrosion resistance, 112.

Upper hainiteformation, 4. 5.morphology, 4-5, 27.properties, 6-7, 27-28. 4J .

Uses, 49. 72, 77. 80. 113, 115.

Welding, 48-49. 76, 78, 113.Work hardening, 21. 43-44, 57. 71-72, 76, 79

105-106. 113.

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