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Cover Front Replica electron micrograph of HP 9-4-n steel after martensitic quenching, showing self-tempered martensiie with cementite particles. Back Thin-foil electron micrograph of 13 Ni (400) maraging steel aged for 4 hours at 900" F (480° C), showing martensite laths; their granular aspect is due to & very fine precipitate, probably a-FoMo. cobalt monograph series cobalt-containing high-strength steels

a critical review of the physical metallurgy of cobalt- containing high-strength steels, and a survey of their processing, 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 Cobalt Rue Royale 66 B-1000 Bruxelles (Belgique).

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

Cobalt Information Centre 7 Rolls Buildings, Fetter Lane London EC4A IJA (England).

Kobalt-Information Elisabethstrasse 14 D-4 Diisseldorf (Deutsdiland). FOREWORD

The first commercial cobalt-containing high-strength steels of the carbide-hardened, maraging and 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 exhibit properties which surprised the industrial world, and perhaps even the scientists responsible for their development. Although the introduction of these steels was obviously preceded by several years of laboratory research, the origin of their outstanding properties remained unexplained for quite a time. For example, in his conclusions to the Journees Internationales des Applications du Cobalt, held in Brussels in 1964, the undersigned, reflecting the general consensus of both speakers and audience, had to admit that the role of cobalt in these steels, and more particularly its 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 that the 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 hardening is 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 martensitic structure; it also exerts a decisive influence on the precipitation kinetics, favouring the formation and 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 the physical metallurgy aspects of these steels was made possible by the availability of examination techniques of greatly increased sensitivity. The importance of these findings, as well as the growing industrial significance of cobalt-containing high-strength steels, prompted the Cobalt Information Centre to devote the third volume in its " Cobalt Monograph Series " to presenting a critical summary of the knowledge on the three groups of steels under consideration. The task of writing the manuscript was entrusted to the Centre 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, properties and uses were also dealt with, though more concisely. This monograph is based, not only on a comprehensive literature survey, but also on the experience acquired at C.R.M. during its long-standing association with C.I.C., as well as on numerous discussions with specialists in this area of metallurgy. The authors gratefully acknowledge the active support received from the following scientists who, in addition to participating in the discussions, also provided initial readings of the manuscript: Dr. J.H. Gross and Dr. S.J. Matas (Chapters III and TV on carbide-strengthened steels); Dr. S. Floreen (Chapters V to VII on Ni-Co-Mo maraging steels); Dr. H. Brandis, Dr. R.L. Caton, Dr. A. Kasak, Dr. A. von den Steinen and Dr. D. Webster (Chapters VIII and IX on stainless maraging steels). They also wish to thank their colleagues both at the Centre de Recherches Metallurgiques and the Centre d'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 proof checking. The authors have attempted to present a coherent but concise review of the vast amount of information available to date on the three families of alloy steels. They believe that this volume will prove useful to scientists concerned with the strengthening mechanisms of high-strength steels, to metallurgists involved in promoting them, and to engineers in their continuing search for new materials capable of meeting the increasingly stringent service conditions imposed by present-day technology. If this monograph stimulates a reader response either to identify those aspects which require further clarification or to suggest ways of improving properties, then progress 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 2i. Solid-Solution Strengthening - 2 I. Phase-Transformation Sirensthening : Bainilic Reactions . 4 2.2.\. Morphology of Bainites 4 2.^.2. Properties o( Bainiies ft 2.2.3. Role of Cobalt 7 2,.;. Phu>e-Transformation Strengthening : Martensiiic Reactions 7 2 .".I. Cieneral Charaaenstics of Martensiiic "(ran>t'ornii:iions 7 I.'.-. Types of Mancnsile 10 :.:O. Morphology of Lath Martensitc 10 I..v4. Morpholoyy of Twinned Manensite II _.?5. Transition from Lath to Twinned Martensite. ... 12 2.}.b. Properties of Martensites 13 2.3.7. Controlled Martensitic Transformation 15 2.-I. Precipitation Strengthening 16 2.4.!. Mechanisms 16 2.4.2. 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 23 3.2.2. Isothermal Transformation Curves 24 3.?. Biiinitic Transformation Structures 24 3.3.1. Bainites Formed on Continuous Cooling .... 24 3.3.2. Mechanical Properties Associated with Continuous-Cooling Bainites 27 3.3.3. Isothermal Bainite in HF 9-4-45 27 3.3.4. Mechanical Properties of Isothermal Bainite in HP 9-4-45 27 3.4. Martensitic Transformation on Quenching 28 3.5. Tempering Reactions 29 3.5.1. HP 9-4-X Steels 29 3.5.2. 5Ni-Cr-Mo Steels 31 3.5.3. lONi-Co-Cr-Mo Steels 32 3.5.4. Retained Austenite 34 3.6. Effect of Alloying Elements on Tempering Response, Strength and Toughness 34 3.6.1. Effect of Carbon 25 3.6.2. Effect of Nickel 35 3.6.3. Effect of Silicon and Manganese 36 3.6.4. Effect of Carbide-Forming Elements 36 3.6.5. Effect of Cobalt 37 3.6.6. 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 , 42 4.2.1. Strength/Toughness Characteristics 42 4.2.2. High- and Low-Temperature Properties 46 4.2.3. Fatigue Behaviour 47 4.2.4. 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 50 5.1.1. Role of Alloying Elements 52 5.1.2. Compositions 53 5.2. Martensitic Transformation 54 5.2.1. Formation and Morphology of Martensite .... 54 5.2.2. Factors Controlling Lath Martensite Formation. . 56 5.3. Ageing of Martensite 57 5.3.1. Precipitation Reactions 57 . 5.3.2. Ordering 60 5.3.3. The Cobalt/Molybdenum Interaction 60 5.3.4. 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 6.1. Primary Processing 68 6.2. Properties 68 6.2.1. Strength/Toughness Characteristics 68 6.2.2. High- and Low-Temperature Properties 72 6.2.3. Fatigue Behaviour .',...' 74 6.2.4. Stress-Corrosion Characteristics 74 6.3. Secondary Processing 76 6.4. Applications 77

7. Ni-Co-Mo MARAGING STEELS — THE ULTRA-HIGH STRENGTH GRADES 77 7.1. Processing 77 7.2. Properties 79 7.2.1. Strength and Toughness 79 7.2.2. High- and Low-Temperature Properties 79 7.2.3. Other Properties 80 7.3. Applications .80

8. STAINLESS MARAGING STEELS — PHYSICAL METALLURGY . . 81 8.1. Background 81 8.2. Effect of Alloying Elements on Equilibrium Structures . . 83 8.3; Transformation Temperatures and Structures 85 8.3.1. Martensitic Transformation 85 8.3.2. Austenite Reversion 88 8.3.3. Retained Austenite 89 8.4. Grain Size 90 5.5. Ageing Reactions 91 ii.5.1. Iron-Chromium System 91 S.5.2. Fe-Cr-Co and Fe-Cr-Ni Systems 92 8.5.-1. Fe-Cr-Co-Mo and More Complex Ailoys ..... 94 8.5.4. Concluding Remarks 98 8.6. Snengih Toughness vs. Structure Relationship 99 5.6.1. General '99 5.6.2. Effect of Retained Ausienite 99 8.6.3. Effect of Delta-Ferrite 100 8.6.4. Effect of Prior Austenite Grain Size 101

9. STAINLESS MARAC.ING STEELS — PROCESSING AND PROPERTIES 101 9.1. Primary Processing 101 9.2. Properties 103 9.2.1. Strength/Toughness Relationship 103 9.2.2. High- and Low-Temperature Properties and Thermal Stability 107 9.2.3. Fatigue Behaviour 109 9.2.4. Corrosion Resistance 110 9.2.5. Stress-Corrosion Characteristics ...... Ill V.2.6. High-Tempemiure Oxidation Resistance 113 9.3. Secondary Processing 113

9.4. Applications 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 the tremendous effort, which has been devoted to the development and understanding o'i this class of materials over the past twenty years. In Figure I, the yield strengths of steels currently used in so-called '• massive structural applications ", on the one hnnr1 and in " specialized structural applications", on-the other, are plotted versus decade from 1850 onwards. This graph illustrates both the progress that has been mac!,- over the vcars with respect to strength of structural steels and the impact exerted by the advent of highly sophisticated constructions such as aircraft or missiles on current ;ind future strength demand, it shows, in particular, a sharp increase in the slope of the curve lor the specialized applications as from about 1950, and the anticipated coming into use of steels with 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 almon immediate notoriousness because of their outstanding properties. The purpose of th s monograph is three-fold : (1) to present as complete a survey as possible of existing and developmental cobalt-containing high-strength steels; (2) to attempt to explain die role 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-day high-strength steels. As in our previous monographs, the emphasis will be put or. the relationship between properties and structure, although more practical data are also provided. 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 of high-strength steels. The subject matter is divided into four main sections. Chapter 2, which is a general review of the hardening mechanisms operative in high-strength steels, is intended to provide the required background for the subsequent sections. Chapters 3 and 4. are devoted to carbide-strengthened steeis, essentially those of the 9Ni-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 the physical-metallurgy viewpoint (Chapter 5), and then from a practical aspect. On account of conspicuous differences in the type and amount of information available on the " conventional " and " ultra-high-strength " grades, it was found convenient to treat these two classes in separate chapters (respectively Chapters 6 and 7).

400 ~~l T 350 Components for engines,, air frames, missiles, . J.30O and other specialized structural application;., S x-250 6 Lreomotives.bridges,' i >^ §200 ships, buildings, TVtowers, —jr— and other massive L ^ ft 150 structural applications g 100 Fig. 1.1.—- Increase in yield strength of steels used in structural applica* SO tions during the ymrs since iS50. 1900 1950 After A.M. 1-ALL [/./]. YEAR

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 contain cobalt and molybdenum, in addition to chromium for stainlessness. Here again, the former chapter 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 the references retained vwis based on criteria of conclusiveness and originality, but the reader >hould appreciate thai this practice has resulted in the omission of numerous papers of \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 gainel at the Centre de Recherche* Metallurgiques during its long-standing involvement in work on 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 mechanisms that are operatise in high-strength steels, and to discuss their influence on the latter's mechanical properties. As in earlier review papers [2.1 to 2.3], the emphasis will be placed on the effect of cobalt on these mechanisms, reference being made in each case to the sroups of steels more particularly concerned. The various classes of steels discussed in this monograph are listed in Table 2.1. which also shows the strengthening mechanisms involved 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 ferrite or austenite. As such, it will exert a solid-solution hardening effect which may be useful in 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 the energy required to move dislocations. Segregation of atoms can occur through interaction between solute atoms and imperfections, or between the solute atoms themselves. Elastic interaction between solute atoms and dislocations leads to Cottrell atmospheres, whereas interaction of electrical origin between these atoms and stacking faults gives rise to the Suzuki effect. Finally,

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

Steels Strengthening Mechanisms Chapter

Low-alloy steels Martensitic transformation (twinned martensite) 2 Precipitation of carbides Cai bide-strengthened martensi'-ic steels Bainitic or inartensitic (lath or twinned) transformation 3,4 (9Ni-4Co. 10Ni-8Co-Cr-Mo) Precipitation of carbides Martensitic transformation (low-carbon lath martensite) 5,6,7 Ni-Co-Mo maraging steels Precipitation of intermetallic compounds

Martensitic transformation (essentially lath martensite) 8,9 Stainless maraging steels Precipitation of carbides and/or intermetallic compounds 2. STRENGTHENING MECHANISMS IN HIGH-STRENGTH STEELS

2 4. 6 ATOMIC % SOLUTE ATOMIC V. SOUJTE 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 ordering reaction in a cobalt-containing solid solution probably contributes to the strengthening of maraging steels on ageing. Some neutron-diffraction data have indicated that hardening of these steels may be due to a low degree of long-range order or a high degree of shon-range order. Data on both lattice dilatation and the changes in elastic constants with solute concentration are required in order to estimate the magnitude of the solid-solution strengthening effect. The changes in the room-temperature lattice parameter, modulus of elasticity and shear modulus of a-Fe brought about by the presence of various solid-solution elements are illustrated in Figure 2.1. It is seen that additipn of up to about 6 at. % Co has a negligible effect on the lattice parameter, whereas both the modulus of elasticity and the shear modulus increase with increasing cobalt additions. Figure 2,2, which is a plot of the size misfit parameter, za = (l/a0) (da/dc), versus the change in shear stress with solute concentration, AT/Ac, shows that little strengthening of a-Fe occurs for low concentrations of 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].

Be. - 5 "g Ptr-30 i*

|9NI.

Fig. 2.2. — Correlation of R.T. solid- solution strengthening of iron-base 002 0.01 006 ODB D.1D 0-12 OK 0.16 0 alloys with size:, misfit parameter. After;^V.G;-;LESLIEJ[2J]/: . i HlliHSIRl NCIIH Sir US

The strengthening effect of elements in substitutional solutions can be influenced, sometimes markedly, by the interstitial content of the solid solution, as well as by more readily controllable variables such as substructure, grain size and temperature [-.7], Work on a O.lC-l2Cr martensitic steel [-.*'] has shown that, for cobalt additions up to approximately 15",,. a hardness incrjmeni of approximately 9 HV per 1 vvt.% Co occurs. The higher solid-solution hardening effect of cobalt in u murtensitic structure can be used to advantage in 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 following section, 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 constituent of steels which is formed by the decomposition of austenite within a temperature range located 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 stabilized austenite. Its morphology changes progressively with the transformation temperature, in

Transformation franl

C-enriched zones Austenite- marlrnsite nodule -•• Acicular ferrile' Nucleation ;ind growth Nudcation and growth Formation of massive bainite showing gran, of upper bainite of lower bainite 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 acicularity of the structure. As illustrated in Figure 2.3, three distinct types of bainite, designated massive, upper, and lower bainites, can be produced in many steels by the athermal transformation of austenite at suitable cooling rates or by isothermal transformation at subcriiical temperatures; the first is also designated granular baiciite (when observed at low magnification) or acicular fcrrite (when observed at high magnification). In the isothermal transformation diagrams for steels with sufficiently high carbon contents, the bainitic field is roughly divided into two horizontal bands corresponding to upper and lower bainite [2.9]. For upper bainite, which forms above 660-750°F (35O-4OO°C) depending on the steel's composition, ferrite plates nucleate within the austenite grain; carbon diffuses av ay and concentrates in the remaining austenite until it finally precipitates as Fe3C carbide between the ferrite plates (Fig. 2.4a). The carbides are often parallel to the axis of the ferrite needles 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 in the same way, but carbon no longer diffuses so readily and cementite, occasionally preceded by e carbides, precipitate within the ferrite plates (Fig. 2.4i.'). The carbides form an angle of approximately 60° with the axis of the ferrite needles. Moreover, it has been shown that the cementite-ferrite orientation relationships in lower baiuite are identical with those that prevail in tempered martensite; examples pertaining to 9Ni-4Co steels will be found in Chapter 3 (Section 3.3.1). Massive bainite forms easily in low-alloy steels during athermal transformation of austenite; it corresponds to microstructures consisting of coarse plates or presenting an almost entirely granular aspect (Fig. 2.5). It occurs over a range of cooling rates that is dependent on the steel's composition, and consists of irregular ferrite grains having a moderately high dislocation density and islands of enriched austenite which may transform either completely or partly io martensite [2.9]. In order to explain both the stabilization of the austenite and the characteristic appearance of massive bainite, it has

1 a) Fully transformed at 930oF (500°C) : b) Fully transformed at.750°F (400°C) : Fig 2i —' Microstructural features of upper bainite. lower bainite. athermal bainitic transformation in 0.1C- lCr-0.5Mo-0.002B sleel fully transformed Fig. 2.4. — Microstructural features of isothermal bainitic transformation at 885°F (475°C). After M.E. BUSH and in low-carbon steel. After K.J. IRVINE [2.12]. <-i-•••":' ''• x 16,000 P.M. KBLLY [2.13]. x 1500 O>H U I i ON i UMNO UK.H-SIRl Nl.ltl Nfl-lUS

been sueuested that dehomogenization of the austenite occurs on holding the latter in us meuistahle temperature range [2.14], Carbon would diffuse under ihe effect of an acmitv gradient, as opposed to a concentration gradien1. towards regions having a high den>i\> of imperfections. As a result, piior to its transformaiion. the austenite would contain a network of carbon-rich regions. From this stage onwards, the theory for the formation of upper bainite can be applied, to a first approximation, to massive bainite (hiah temperature of formation, low cooling rate) : nucleation of supersaturated ferrite platelets within a carbon-depleted region, followed by their growth and the possible rejection of the carbon, towards tht untransformed austenite. The fairly large size of the hainuic ferme areas in massive bainite results from the more rapid growth of the platelets due to the prior dehomogenization of the austenite. The austenite islands are either swallowed up during growth or imprisoned between two ferritic regions. Their stabilization can 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; the dislocation densit) can also be a stabilizing factor.

2.2.2. Properties of Bainites Both the isothermal and continuous-cooling (atherma!) b.iinites have been investigated extensively 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 greater commercial importance and have consequently formed the subject of m;;ny of the recent investigations concerned with the microstructure / mechanical property relationship in bainitic 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 contributions from the strength of pure iron i.nd the effects of alloying elements in solid solution, the P'ior austenite grain size, iiod the dislocation density. The strength is raiser1 above the base level by the presence of the islands referred to in the preceding section, through a two-phase effect similar to that which leads to the strengthening of fetnte-pearlite aggregates by pearlite colonies. Isothermal bainites §we their base strength to the same factors as massive bainites but with an increased dislocation-density contribution, the acicular ferrite matrix being rich in defects. The strength is raised above the base level by further contributions from the dislocation density, which increase as the transformation temperature decreases. The carbide precipitates also exert a large strengthening effect which is dependent on both the number and shape of the carbides present. As regards toughness, it is known that, for numerous steels, lower bainites are tougher than upper bainites. Since the microstructure of the steels is controlled by the transformation temperature, a relationship exists between strength and transformation temperature. This is illustrated in Figure 2.6. The diagram shown in Figure 2.7, on the other hand, indicates the combination of yield strength and impact transition temperature which can be obtained in ferritic-pearlitic, bainitic, and quenched-and-tempered low-carbon steels. The microstructure of the latter consists of acicular fernte together with small precipitated carbides which 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 strengthening factors in these three types of microstructure. The base line in this diagram represents a relationship between yield strength and grain size. For the C-Mn-Nb steels, strengthening beyond that related to grain size is due to carbide precipitation. In the case of bainitic steels, 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 VELD. STREN5TH,'MN/m* „. too son _6QO TOO eoa

• MAR- . BAINITE M FERRITE] TENSIIE! 1 IPEARUTEI200

.5 . 10 15 20 -iooo..noo wuo ••~KT.Tr:.:..- ~~ ./•"- 80; GRAIN SIZE. d''4 mm''4 Temperature dt maximum rate of lransformailon,.*F YIELD STRENGTH; »> psi.

Fig. 2.6. — Relationship between Fig. 2.7. — Relationship between yield Fig. 2.8. — Relationship between grain transformation temperature and strength and impact transition tempera- size and yield strength of low-carbon strength of low-carbon bainitic ture of low-carbon steels for different steels for different microsiructures. steels. After K.J. IRVINE [2.12\. microstructures. After KJ. IRVINE [2.12]. 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 different from 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 solution in the martensite; hence, the block representing the martensitic structures lies just above the 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 scarcely influences their hardenability and grain size, and does not exert any appreciable solid- solution hardening effect on ferrite. At most, when present in large amounts together with other alloying elements, cobalt might promote strengthening by leading to a better dispersion of the carbides, or to the formation of intermetallic compounds. Also, in some low-alloy high-strength steels, addition of small amounts of cobalt may decrease the brittleness [2.5]. However, there is an important group of cobalt-containing steels in which the bainitic transformation is used to advantage, viz. the 9Ni-4Co steels, which are characterized 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 a review of the general features of the martensitic reactions. A fuller treatment of the subject will 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 low temperatures and this leads to the production of a new phase, martensite. As opposed to a conventional nucleation-and-growth reaction, the martensitic transformation is characterized by the fact that it cannot be suppressed by rapid cooling, even using very high cooling rates (up to 32,500°F/s, i.e. 18,000°C/s). In addition, it requires no diffusion . OHALI'-l OM MMMi HK iHSl Kl\(. I H STtllS

or tiuermixinsi of the atoms involved in the phase change; the transformation products inherit ihe composition of the parent phase and each atom lends to retain iis original neighbours. Finally, martensiiie reactions are displaeive or shear-like, in that the atoms mo\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 associated with the transformation from the parent to the product phase, which reaches as much as 10",, in ma,lv commercial alloys. An experimental indication of these transformation strains 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 result of the prior thermal and mechanical history of the material but is generally iiniepcmlcni ni tin- («<>/i"ji!j rate. For highly alloyed ferrous materials, the martensitic transformation tray also be induced at temperatures above the A/« point by plastic deformation; the highest 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 primarily while the temperature drops below .W», giving rise to the well-known atlwrmal type of martensitic reaction. However, there are instances in which ivjcleation may occur isu- ihermalh: and the martensitic plaits will then exhibit rapid growth even at a fixed temperature. (iii) Both the >ize of the martensiie grains and the A/.< temperature are dependent on the ausienite grain si/e, which itself is a function of the time and temperature of the austenizing treatment : 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 rate or holding the alloy within a specific temperature range. Stabilization of austenite in carbon steels is dependent on this characteristic. (v) A temperature hysteresis exists between the heating and cooling reactions; the temperature at which marlensite starts to transform to austenite on heating at the rate of 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 most acceptable to date is that which proposes that embryos of the manensitic phase form within the austenite. and that some of these reach a critical size which enables them to grow at temperatures beiow Ms [2.21]: the small regions of martensite thus formed are separated from the austenite phase by a semi-coherent interface. A recent structural investigation of the 18Ni(3OO) maraging steel [2.22] did in fact show that, on cooling below A/«. martensite platelets nucleated within the austenite grains themselves and not preferentially at the grain boundaries; on the other hand, their growth was stopped either by the prior grain boundaries or by previously formed platelets. Moreover, the platelets did not appear to nucleate and grow uniformly within the structure, since some regions were almost completely transformed to martensite, while others were still entirely austenitic.

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 resulting mechanical properties of the steel. In low-carbon, low-alloy steels, the M, temperature lies between 570 and 750 F (300 ;ind 400 C) and full transformation can be obtained on cooling through the martensite transformation range to room temperature. When appreciable amounts of alloying elements are present,-however, the Ms temperature is depressed to near room temperature and the M, temperature, i.e., that at which the martensitic 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 15 Co 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 addition of up to 8l,,Co to an Fe-22.5Ni alloy raises the M* temperature. This effeci applies to Fe-Ni alloys and maraging steels in general; it may result from the lowering of the austenite shear modulus by the cobalt addition [2.29]. However, a direct relationship between Ms and cobalt content is not always observed. For instance, the Xis temperature of Fe-18Cr and Fe-12Cr-4Ni steels decreases with increasing cobalt contents. Figure 2.9 also shows that increasing the cobalt content of Fe - 17 to 19Ni - L5Mo alloys from 15 to 19.5% lowers their A/* temperature; this is in contrast to the behaviour observed in Fe-Ni alloys and maraging steels for cobalt contents below 8%. Data on the effect of other elements on the Ms temperature of steels are given in Chapters 5 and 8 (Sections 5.2.2 and 8.3.1).

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

As regards the As temperature, it is strongly depressed by addition of appreciable quantities of alloying elements such as nickel; this has a direct bearing on the choice ot the tempering temperature. Cobalt also has a depressive effect, though much weaker than that of nickel1 this effect is retained in chromium steels (Fig. 2.9), but in Fe-Ni alloys and maraging steels in general, cobalt has been shown to raise the />„ temperature. tOlUl |-u»\l VIM Mi Hll.il s( RIA(,1H SIM IS

Fig. ; 10 • Microsiructural features if lalh manensiie in Fe-1.94Mo alloy. X 14,000 After 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 these two martensitic products, which differ with respect to composition, range of formation in a given alloy system, crystallography, morphology, and fine structure, has led to the use of a multiplicity of terms to differentiate between them. A definite preference for the term " lath martensite " for one form and " plate martensite " for the other was recently reported [2.17]. Although the former will be used in this monograph, an alternate designation, 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, which are planar and lie along one direction, are generally aligned parallel to one another in groups that have been termed packets, fragments, blocks or sheaves. These packets are the predominant feature of the microstructure and the individual martensitic laths are visible as a fine substructure within the packets. Several packets are found within a prior austenite grain. The widths of the units which make up a packet of martensite range from less than one tenth to several microns, with the most frequently occurring width being between 0.1 and 0.2 \xm. Adjacent martensite laths may be separated by low- or high-angle boundaries or may be twin-related [2.30, 2.31], but this does not appear to be generally true [2.32]. The orientation relationship between martensite packets and the austenite parent phase are generally of the Kurdjumov-Sachs type, i.e. :

(ii i).,//(oil)* Uio]T//rjTi], The interface between martensite packets and the austenite is parallel to a jlll)rtype plane [2.32]. Moreover, the orientation relationships between adjacent laths forming the same martensite packet are generally of the type :

(1 ">)»,//(110)«2 [TllW/[001]a2 Lath martensite is also characterized by its substructure which consists predominantly of a high density of tangled dislocations within the laths. It has been stated that dislocations in structures of this kind tend to lie in <11 l>a directions and are predominantly screw in nature [2.33].

10 2 STRI- MECHANISMS IN HIGH-STRENGTH STEELS

The lath martensitic transformation differs from the twinned manensitic transformation in carbon steels mainly in that the martensite packets grow more slowly and the changes in the austenite-martensiie interface with time are not the same. The latter involve a very short-range diffusion process, over two 10 three atomic layers, in order to relieve the 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 to 25 "„ Ni is of particular importance within the framework of this monograph since it is largely responsible for the properties of the maraging steels to be discussed in Chapters 5 lo 7. The high density and quasi-uniform distribution of dislocation.* in this structure favour a belter age-hardening response (see Section 2.4) by providing a large number of nuclcation sites and increasing the diffusion rales, hence ensuring a more uniform distribution of finer precipitates.

2.3.4. Morphology of Twinned Martensite Twinned martensite differs from lath martensitc in that adjacent plates do not form parallel to one another. The first, plates formed lend to span their parent austenite grains and effectively partition the austenite, thus limiting the size of subsequent plates, as shown in Figure 2.1 la. The effect of this partitioning is to produce a wide range of plate sizes in 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 recently the 11 IOjM type was observed for an Fe-1.82C alloy [2.35]. Frequently a fine dislocation structure consisting of parallel arrays of screw dislocations is found to coexist with fine transformation twins in plates of Fe-Ni martensite [2.33]. In this monograph, the twinned martensitic transformation is of interest in the discussion of carbide-strengthened martensitic (Chapter 3) and stainless maraging (Chapter 8) steels.

a) in Fe-1.86C alloy, showing auto- b) in Fe-32Ni alloy, showing twinned catalylically formed plates. * 440 substructure of plates. x 40,000 Fig. 2.11. — Microstructural features of twinned marlensiie. After G. KRAUSS and A.R. MARDER [2.17].

11 HU I i i:\l \l\l\i. I 111. II SIU! Null! SUMS

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

Several explanations for the transition from lulh to twinned nuiriensite have been proposed The occurrence of one or the mher type is controlled. ;il least partly, by the carbon content of the -steel: as shown in r is;ure 2.12. the martensue of plain-carbon steels containing up 10 U.4",,C consists mainly of lath nutrtensile. but the ratio of lath to twinned nianenOie decreases rapidl> for higher rarhon conteills. On the other hand, the morphological charnic which develops in Ke-C alloys with increasing carbon content h.is been aitrihiiied to twin formation below a critical A/, range of 570 to 425 1- tU)0 to 22II O \2 .U]. Another stud} has led to the conclusion that not oniy temperature but al-o allo\ composition may signiticantly intluence the extent of twinning in the marlensite I-Vi: the important factor appears to he the relative magnitudes of the critical resolved shear -tresses (or twinning and slip at a giv.'n temperature for a particular alloy compx-iiion It h.is ,i!>o been suggested that low slacking-fault energy (SI lit favours the formation of lath nuirtensite [2Jx]. However, there are addition elements which have opposite effects on the SH of austeuile. yet favour the formation of twinned martensite when added in increasing amounts [-./"]: this >- ''he ease for manganese, which lowers the SFE. and nickel, which raises it. According to a fourth proposal [-J-J. lath martensiie in iron alloys 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 the imerstitiul atoms as the martensite forms. The major exception to this hypothesis is lhat cubic 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-C steels [2.3'J and ?.4<>. resp.]. it was stated that the V, temperature is not a sufficient indication of the extent of internal twinning in martensite when comparing two different steels. Howe\er. for a given steel, the temperature range over which martensiie forms determines the extent of twinning in the martensite units. It was suggested that the driving force for transformation, as given by the area of the transformation hysteresis loop, may give a better indication of the occurrence of twinning in martensite. A larger hysteresis gap would mean a greater driving force and consequently a greater chance for the martensite units 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 corresponding phase diagrams, it was also suggested that the y-stabilizcd systems form lath martensite and. with sufficient additions, twinned or h.c.p. (e) martensite as well, whereas the y-loop systems can transform to laih martensiie only [2.17]. This proposal lends support to the temperature hypothesis mentioned at the beginning of this section. The limited austenite solid solution in the y-loop systems does not permit sufficient alloy additions Ki reduce the V/« temperature of the austenite to the critical temperature. On the other hand, sufficient alloy additions can be made in the case of the y-stabilized systems to promote twinned martensite formation. The effect of cobalt on iron alloys does not contradict the y-stabilization and hence the temperature 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-/, temperature so 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-C steels [2.39, 2.40] showed that cobalt is not effective in retaining lath martensite for increasing carbon contents. Although the addition of cobalt raises the MH temperature of such steels, it does not reduce twinning unless the Mx temperature is such lhat the critical resolved shear stress for slip is lower than that for twinning.

12 1. STRENGTHENING MECHANISMS IN HIGH-STRENGTH STEELS

NICKEL, at.% ICO 2 < 6 8 10 _B 20 25 30 35 RANDOM _ i LATH ! TWINNED

RELATIVE VOLUME % LATH MARTENSITE I

wt.%C

Fig. 2.12. — Effect of carbon content on relative volume frac- Fig. 2.13. — Separation of solid-sol- tions of lath and twinned martensites in Fe-C alloys. The ution hardening and transformation AA, temperatures and volume percentages of retained austenite substructure effects in Fe-Ni alloys. are also shown. After G.R. SPEICH and W.C. LESLIE [2.36]. 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 strength arise from : (i) ultrafine subdivision of the product-phase grain into twins; (ii> significant increase in the dislocation density; (iii) re^.nement of the parent grain size in the absence of twinned products; (iv) additional solid-solution hardening; (v) additional hardening through ordering of the product phase. However, martensites are not necessarily hard and strong, which means that the mere occurrence of a martensitic transformation does not per se confer conspicuous strength upon the product of the reaction. Even in the case of steel, the martensitic transformation does not result in high strength if the carbon content is low. In the absence of interstitial solutes, the greatest hardness that can be attained by a martensitic transformation in unalloyed iron or dilute iron-base substitutional solid solutions is about 260 HV, with a yield strength of about 100,000 psi (700 MN/m-). This strength is little affected by significant changes in the structure, such as the transition from lath to tinned martensite [2.31, 2.32]. In Figure 2.13 the yield strength of quenched and of .-ecrystallized Fe-;Ni alloys is plotted against the square root of the nickel content; the flow stress c^-of the quenched alloys can be separated into three components : af = (<*a + 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 lath boundaries, and the cell walls or internal twins within the transformed regions). It is seen that the solid-solution hardening effect of nickel accounts for about three-quarters of the overall strength of Fe-Ni lath martensite containing 20%Ni. Figure 2.13 also shows that changes in

13 Oil U I" l 1>\1 UMNt: Hll.H-SIKlAtiill S1T11S

I it;. 2.14. — Transl'ormalinn sub- SII'UL'UIIV strengthening cll'ocis in 1 e-C'o alloys. After AC. 4S0 Ssuiv and I:..K. linv |.\V.>'|

400:

350^

300S Annealed Ferrite_

2S05

nk 76 73 SO Co CONTENT at.% structure changes from a lath- to a twinned-martensile type. As regards Fe-Co binary alloys. Figure 2.14 shows that progressive strengthening occurs as the cobalt content is increased, culminating in a 44",, increase in O.I "„ offset yield strength due to the formation of a lath mariensite substructure. In Fe-C iillo\s with low carbon contents, up to about 0.05",,. most of the strength rr the mariensite seems to deri\e from the high density of dislocations, which is probably about 10i; to 10*; per cm:. In alloys with higher carbon contents, it seems likely that the strengthening due to carbon exceeds by far that due to the substructure [2.44]. Recent work on Fe-Ni-C alloys [2.45] has shown that, at carbon levels higher tl.an 0.3 wt.",,, the lath (cubici martensites are significantly stronger because they deform by slip whereas twinned I tetragonal I martensites with the same carbon content deform by twinning. It was suggested that deformation by twinning is suppressed in the high-carbon cubic martensius because in these structures all twin variants inevitably carry a large fraction of the carbon atoms into high-energy, non-octahedral sites. The work-hardening behaviour of lath martensites. the temperature- and strain-rate dependences of their flow stress, and the derived values of the activation volumes and energies for deformation, are all reasonably consistent with the behaviour observed in iron or low-alloy steels [2.46]. Furthermore, their Pelch slope values, based on the prior austeniie 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 of cold-worked iron, as was proposed earlier [2.47]. One striking difference in properties between lath mariensiie and cold-worked iron is that the ductility of the former is reasonably good, as is exemplified by the Fp-l8Ni martensite even when tested in liquid hydrogen. Studies of the plastic deformation of an 18",,Ni binary [2.48] and several Fe-Ni-Cr ternary [2.49] lath martensites showed that the alloys all behave in a rather similar fashion. In the Fe-Ni-Cr ternary alloys, higher nickel levels considerably improved the impact properties at sub-zero temperatures Chromium was also slightly beneficial to toughness. 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 the fact that the number of mobile dislocations is higher in lath martensites [2.48]. The availability of mobile dislocations in the structure would, in general, minimize cleavage and allow slip to take place. Thus, with regard to the maraging steels, one obvious advantage of the Fe-Ni lath martensite matrix is that the flow stress is considerably improved 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 fairly large increases in the elastic limit, and small increases in the yield and ultimate tensile

14 2. STRENGTHENING MECHANISMS IN HIGH-STRENGTH STEELS

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-Cr alloys, the increases in the elastic limit were inversely related to the A/« temperature of the alloys. 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 the martensitic transformation is primarily responsible [2.49]. There is also some evidence that, after tempering, lath martensite gives belter toughness than twinned martensite. The role of microtwins in lowering the toughness of Fe-Ni-Co-C steels was demonstrated by comparing heavily twinned with untwinned martensile, and twinned martensitc with bainite [2.J9]. At similar strength levels, the toughness of lower bainite was found to be superior to that of heavily twinned martensite, but inferior to that of untwinned martensite. Since, as stated • • Section 2.3.5, cobalt is not effective in reducing twinning in the higher-carbon steel* f :his type, the latter's toughness is not enhanced by cobalt additions; in fact, add.:ii<.;.i i t more than 4%Co was found to be detrimental to the toughness of high-carbon .•.•.•.. 2.39]. In conclusion, this section on mechanical properties again emphasizes the importance of 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 the Fe-Ni base. From this point of view cobalt, which raises the Ms temperature of the base alloy, is very helpful in that it allows larger amounts of alloying elements to be added while ensuring that the lath martensitic transformation remains possible. In addition, although cobalt does not affect the martensitic hardenability, it may possibly be used to increase 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 martensite transformation temperature range is carefully controlled so that it is close to room temperature. The microstructure is also controlled so that it contains a small percentage of 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 cold working after cooling; the martensitic transformation is induced either by primary tempering which, through precipitation reactions, raises the Ms temperature of the residual austenite, or else by refrigeration. The fact that the microstruciure is controlled means that the chemical composition is very critical, and consequently these steels are difficult to produce. Cobalt is a particularly useful element in such steels because it reduces the stability range of S ferrite without depressing the Ms temperature and it car. be used to balance 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 develop such steels. It is just on the borderline for S-ferrite formation and its martensitic transformation range lies just abOVe room temperature. I" order to modify the tempering characteristics, alloying elements such as molybdenum (one of the useful carbide-forming elements) and copper or aluminium (which are useful age-hardening elements) are added. When molybdenum is used, its ferrite-forming tendency is conveniently balanced by cobalt;'simultaneously, the chromium and nickel contents are reduced to allow addition of the increased amounts of cobalt and molybdenum. A typical composition of this type is 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

I" - • — ~—

"•"—— " -. 2% Co . 0% ci —_ h// t as 1 1 6. 8 JO 20 <*) 60 80100 350 AGEING TIME, hours 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 age hardenability of the modified steels. This is shown in Figure 2.15 for a Cr-Ni-Mo and a Cr-Ni-Cu steel. Since the general sequence of the tempering changes is identical to that in Co-free steels, it has bean suggested that the effect of cobalt is almost entirely due to solid-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 the tempering process (Fig. 2.16). The first, stress relief, corresponds to precipitation of e carbide at temperatures of the order of 200-4003F (100-200°C); in actual fact, it is preceded by carbon segregation or pre-precipitation clustering [2.36]. The second stage, which takes place between 400 and 600°F (200 and 300°C), involves the decomposition of retained austenke. Precipitation of cementite (FesC) in most carbon steels tempered between 500 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 progressively transforms to spheroidal precipitation. Substitution of more stable alloy carbides for cementite 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. Before discussing the carbide and intermetallic-compound precipitation reactions in greater detail, the mechanisms which account for the strengthening role of precipitates will first be 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 precipitates and ordinary precipitates. The former may be regarded as a compositional change in the solid solution, with or without a structural change. If there is a structural change, then it is a perturbation of the structure of the solid solution. An ordinary precipitate may be coherent, semi-coherent, or incoherent with the matrix. Actually all types of precipitate

16 2. STRENGTHENING MECHANISMS IN HIGH-STRENGTH STEELS

TEMPERING TEMPERATURE'C -59- • i tj formation and growth • of Intennttoilic compounds^, Singe V Formation ana* growth of alloy carbldtt^i Stage IV low-carbon marwmltt + £ -r-ferrlte + cementite | Swge III

Stoge II

Stage 1

w in m m TEMPERING TEMPERATURE'F

Fig. 2.16.- - 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 of coherency are more effective in this respect. The problem which must be considered in order to understand precipitation hardening is 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 and Lenel's model). In this model detectable plastic flow would occur only when the particles are being sheared or broken by the passing dislocations. Dislocations pile up against second-phase particles and the latter rupture whenever the accumulated stress is large enough. Thus, according to the final form of Ansell and Lenel's model [2.56], the yield strength of a dispersion-hardened material is given by :

7 P* (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 each particle (Orowan mechanism). In this model, the plastic strain results from the expansion of the dislocation loops surrounding the particles which intersect the glide plane. The initial shear yield strength is given by the modified Orowan relationship [2.54] : . Gb 2 In with q> = — (2.3) 2b

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 the particles and the mean planar interparticle spacing, and v is Poisson's ratio. (iii) The dislocation line bypasses the particle by cross slip, leaving dislocation segments behind (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 been proposed to account for the yield stress of precipitation-hardenable alloys [2.55, 2.56].

17 COBAU-l/ONl AIMNti IlKiHSTRFNGTH STKRS

TEMPERATURE.t TEMPEHINS TEMPERMURE, »C 300 400 530 100 200 300 400 500 600

800-

600 BOO 600 800 1000 1200 TEMPERATURE.*F quenched TEMPERING TEMPERWURE,''F.

Fig. 2.17.— Effect of tempering temperature (time: Fig. 2.18. — Effect of a 3.3% cobalt addition on I hour) on hardness of Fe-C marlensites. the tempering resistance of an AISI 4340-type steel. After G.R. SPEICH and W.C. LESLIE [2.36]. 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 alloy carbides during heat treatment, which can produce marked secondary hardening. Figure 2.17 summarizes the complete process of tempering in Fe-C martensites and the corresponding hardness variations. Alloying elements such as molybdenum and vanadium are widely used because of the beneficial effects obtained by precipitation of their alloy carbides. Cobalt does not form carbides but may affect carbide precipitation indirectly by preventing recovery of the dislocation substructure during tempering; this provides more nucleation sites and a finer dispersion of the dislocation-nucleated carbides. These features 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 iron containing 2 to 5 wt. % Mo and about 0.2 wt. % N, GP zones have been shown to develop at 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 a partly coherent intermediate metastable phase, TJ' ; (iii) precipitation of a stable incoherent phase, TrFe.,Mo3N. The material thus produced is very hard, the yield strength of the iron reaching about one-half of its theoretical maximum. Moreover, the GP zones are unusually stable; they persist, and the material remains hard, even after heat treating for several hours at temperatures approaching 1290T (700°C). The production of GP zones in 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-interstitial solute-aiom zones must precede the homogeneous precipitation of alloy-element nitrides and cait-ides in most systems. Normal quenching plus ageing treatments promote heterogeneous 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 the relevant chapters. As regards high-strength low-alloy steels, they are used in the martensitic condition, after some light tempering to obtain adequate ductility. In actual fact, tempering is 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 which accompanies this treatment, use can be made of a non-carbide-forming element such as silicon, which has quite a marked effect in stabilizing the precipitates, and hence in improving the tempering resistance of the steels. Cobalt, even in small amounts, seems to 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 would be to impede the growth and coalescence of the precipitated particles by concentrating in the surrounding matrix, as is the case in high-speed steels. This effect has been utilized in commercial steels. Figure 2.18 shows that addition of 3.3 %Co to an AISI 4340-type alloy steel (0.4C-0.6Mn-0.3Si-0.8Cr-0.25Mo-1.8Ni) has a significant influence on the tempering 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 the martensite by shifting FeiC carbide formation to higher temperatures [2.60]. This steel has 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 the formation of intermetallic compounds, is used in numerous steels including the maraging types. 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 (see Chapters 8 and 9). Both families of steels contain fairly large cobalt and molybdenum additions. Depending on the composition of the steel, hardening results from the precipitation of both carbides and intermetallic compounds, or of intermetallic compounds alone. After quenching, ageing at a relatively low temperature leads to a sharp hardness increase, 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 operative on ageing Fe-Ni-(Co)-Mo maraging steels. The first takes place within the martensitic matrix and is predominant for ageing temperatures below 840°F (450°C), while the second occurs preferentially on dislocations and is predominant at higher ageing temperatures. The first mainly involves the formation of coherent ordered precipitates, whereas the second leads to the precipitation of intermetallic compounds which have been identified as 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 of metastable 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. Several investigators have pointed out that precipitation strengthening of the magnitude observed for maraging steels can be accounted for quite reasonably by Orowan's model involving dislocation 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 distributed in the matrix, with an average interparticle spacing of about 300-500 A; the corresponding strength increment derived using Eqn. 2.3 is of the above-mentioned magnitude. Further- UHM \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 instances of dislocations bowing between the particles. Finally, the initially high work-hardening rate of the steels is. in general, more consistent with the Orowan mechanism. On the other hand, some authors [2.6-i, 2.65] have suggested that Ansell and Lenel's model might be more plausible. In this case, the high strength in maraging steels would be due noi only to the very fine interparticle spacing, but also to the very high shear strength of the precipitate 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 show that the effective shear stress varies as thai of mild steel and that the major strengthening is due to long-range internal stresses such as would be developed by fine precipitate particles. One of the most interesting strengthening effects in maraging steels is that due to the combination of cobalt and molybdenum [2.46]. It has been shown that the age hardening obtained in Fe-Ni. Fe-Cr and Fe-Ni-Cr steels when cobalt and molybdenum are present together is much greater than the sum of the strength increments produced when these elements are added individually. The role played by cobalt is not clear, since this element does not significantly participate in the precipitation reactions; numerous explanations have 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 to strengthening by some of the mechanisms operative in all metals : solid-solution strengthening, grain-boundary and interface effects, dispersion strengthening, and strain hardening. The influence of TMT is particularly strong in steels, where the phase change from austenite to ferrite (or martensite) on cooling and the presence of carbon, which has a greater solubility in the aus'.enite and therefore precipitates in the ferrite, combine to maximize the effects. In recent review papers, TMT's were classified according to the position of the deformation in the heat-treitment cycle, i.e.. the structure that is deformed and the final structure

STABLE AUSTENITE

Critics! Temperature

V Start of Isothermal Transformation

Room Temperature

Fig. 2.19.—Schematic time-temperature-transformation diagram showing thermomechanica! treatments. After E.B. KULA [2.67].

20 2. STRENGTHENING MECHANISMS IN HIGH-STRENGTH STEELS

that is formed. These treatments are shown in Figure 2.19 on a conventional isothermal transformation diagram for steel. The TMT classification is then as follows : Class I : the austenitc is deformed before transformation and the martensite forms from this strain-hardened austenite. Ausforming, ausworking, ausrolling and hot-cold working all 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 forms during the deformation of the steel in a metastable condition. Most frequently this treatment is applied to stainless steels, which have MK temperatures slightly below room temperature. Class III : the deformation is carried out after the auslenite transformation and may be followed by reageing. These treatments correspond to the strain ageing of austenite transformation products (martensite, tempered martensite, bainite or pearlite). Various terms have been coined to designate such treatments : flow tempering, strain tempering, marstraining, and tempforming. The above classification is useful because within one group similar strengthening mechanisms are operative. The Class II TMT's are probably the least complex. In the low-carbon metastable austenitic stainless steels to which this treatment is generally applied, the strength 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 refinement of the austenite and the resulting martensite, and also by the presence of defects in the martensite, which have been inherited from the strain-hardened austenite through the phase transformation. The major part of the strength improvement is associated with the 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 during subsequent reageing. The carbides in the martensite increase the rate of work hardening of the martensite and hence the dislocation density. During reageing, some dissolution of the carbides may occur, leading to Cottrell locking and stress-induced ordering around the dislocations, and ultimately to a reprecipitation process yielding a finer carbide dispersion. It was also found that plastic straining after initial ageing, followed by a further ageing treatment, can produce marked changes in structure and properties; in particular, prior plastic straining can considerably shorten the time cycle for heat treatment [2.68].

Many investigations have been made of the effects of TMT on carbide-strengthened 9Ni-4Co steels (Chapter 4), Ni-Co-Mo maraging steels (Chapters 6 and 7) and stainless maraging steels (Chapters 8 and 9). The strengthening response of the 9Ni-4Co steels to TMT Was evaluated with particular emphasis on fracture toughness (c/. Section 4.2.1) : as regards the conventionally heat-treated steels, both strain tampering and ausforming of martensite extend the strength range of these steels and produce about the same toughness 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 while retaining acceptable toughness. As regards 18%Ni maraging steels, most of the work has 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 improvements in the'final strength; on the other hand, marstraining gave more substantial strength increases. Ausforming and strain-ageing treatments have also been used on AFC-77, a cobalt-containing high-strength stainless steel, in order to improve the compromise between strength and toughness (cf. Sections 8.4 and 9.2.1).

21 (.'OHM 1-uAI \|\l*-(, HK.II SI Rt NtilH KUIHS

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 toughness u and adequate weldubility. The HP 9-4-X sieels contain *> -oNi. 4°0Co, small amounts of the carbide-forming elements chromium, molybdenum and vanadium, and carbon in contents ranging from 0.20 to 0.45 "o. depending on the grade. They were developed by 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 States Steel Corporation with the purpose of providing a structural steel for the fabrication of large pressure \essels and hydrospace vehicles which incorporate weldments of heavy sections. Steels of this type were initially considered for use as filler metals capable of pro\iding yield strengths of 170,000-200.000 psi (1200-1400 MN/m-) combined with satisfactory 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 only to 10°uNi grades, but also to some containing 5",,Ni [3.6. 3.7], was based on the concept of combining the advantages of standard carbon martensite with those of maraging in compositions designed to exhibit a " dual strengthening " effect [3.8]. However, later investigations [3.V] showed thai strengthening is entirely due to carbide precipitation during tempering, though according to recent, as yet unpublished, work [3.10] the contribution of intermetallic-compound precipitation to strengthening would be significant for material in the overaged condition. The microstruciural Features of these sieels will be examined in the following sections, as well as the way in which they are affected by heat treatment; their relationship with the basic mechanical properties will also be reviewed. Although some of the steels listed in Table 3.1 are no longer commercially available (this is the case for HP 9-4-45, superseded by certain lower-alloy steels, and HP 9-4-25, discarded in favour of HP 9-4-30 or HP 9-4-20), they will not be excluded from this and the following chapter, in view of the valuable information which their development and study has generated.

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

Steel designation C Mn Si P S Ni Co Mo Cr V Al Year

0.42- 0.10- 0.10 0.01 0.01 7.0- 3.5- 0.2- 0.2- 0.06- HP 9-4-45* — 1962 0.48 0.25 max. max. max. 8.5 4.5 0.35 0.35 0.12 0.29- 0.15- 0.10 0.01 0.01 7.0- 4.25- 0.9- 0.9- 0.06- HP 9-4-30 — 1962 0.34 0.35 max. max. max. 8.0 4.75 1.1 1.1 0.12 HP 9-4-25 * 0.24- 0.10- OJO 0.01 0.01 7.5- 3.5- 0.35- 0.35- 0.06- 1962 0.30 0.35 max. max. max. 9.0 4.5 0.55 0.55 0.12 — 0.16- HP 9 4-20 0.10- 0.20 0.01 0.01 8.5- 4.25- 0.9- 0.65- 0.06- 1966 0.23 0.35 max. max. max. 9.5 4.75 1.1 0.85 0.12 — lONi-Co-Cr-Mo 0.12 0.004 0.11 0.001 0.006 10.0 7.8 2.0 1.0 — 0.02 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 steels of 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-9-4 are 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 which starts at about 445 F (230 C), a broad bainitic transformation range beginning at about 930 F (500 C), with a critical cooling rate of about 270 F/min (150 C min). No transformation to pro-bainitic ferrite was observed, even at the slowest cooling rate used (about 4.5'F/min, i.e.. 2.5 C/min). On decreasing the cooling rate, a considerable, though progressive decrease in hardness is evident, and the residual austenite content is found to increase 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 cobalt raises the transformation points, both on heating (Af and Af) and on cooling (Ms and = fls); 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 nucleation rates. It is also apparent that the residual austenite content in the bainitic range is lower for the 9Ni-4Co than for the 9Ni steel. The reason for the difference between the 9Ni-4Co laboratory heats and the industrial steels studied lies essentially in the latter having a lower nickel content (7.5 "„) and containing the carbide-forming elements chromium, molybdenum and vanadium. The diagram

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

600

ir 200 '° 51 I IB?

• + PHUF l-M-'t !''-- 4TTTffli

» 10 U TIME 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]. Steel 0.3C-9Ni (2-9) 0.3C-9Ni-4Co (Z-9-4) HP 9-4-20 HP 9-4-30 Austenizing temperature 15IO°F(82O°C) 1455°F (790°C) 1525°F (S30°C) I525°F(830°C) Time at austenizing temnerature 30 min t5 min 30 min 30 min A, 1204°F (651°C) 1330oF(722°C) 1567°F (853°C) 1546°F (841°C) A, 1114oF(602°C) 1124°F (607°C) 1198°F (648°C) 1231°F (666°C)

23 Ui|l\ll-iOM MMSi. llliiliMRlMillI MMIS

1 10' K>2 TIME.minutes

Fin. 3.2. — Isothermal transformation diagrams for HP9-4-20 after R.T. AULT [3.12]. HP9-4-25 after 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 the critical cooling rats Co 55 Fmin. i.e., 30:C/min) and 6.,-. On the other hand. A/* is only slightly raised, the effect of adding the carbide-forming elements being more than counterbalanced by that of decreasing the nickel content. Finally, the cooling rate does not affect either the hardness or the residual austenite content: the latter, which is of the order of 5",,. 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 increase in the \tx and Bs temperatures, as well as in the critical cooling rate, which is raised to 600 F,min (33OC. min). The residua! austenite content increases on passing from the martensiiic 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. These diagrams 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-tempered condition. 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 combination of strength and toughness, which is not the case for the lower-carbon grades or the 10Ni-8Co-Cr-Mo steel. Furthermore, bainitic structures may develop in heavy-section weldments which are not subjected to post-welding heat treatment. It is thus appropriate to discuss these structures here; this will be done essentially in relation to the HP 9-4-X steels. The bainite structures formed both on continuous cooling and isothermal holding will 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 transformation 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 of large areas of bainitic ferrite and white, less deeply etched zones composed of an intimate mixture of martensite and residual austenite. Although the martensite is clearly visible,

24 3. CARBIDE-STRENGTHENED STEELS — PHYSICAL METALLURGY

:^ rci

.^ -v «y

4\ + ** \

*>• >,.

lMtun 3000 Fig. 3.3. — Replica electron micrographs of experimental 9%Ni steels after cooling 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 that cobalt 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 differ considerably from those of the steels without carbide-forming elements. Following slow cooling 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 fine carbide precipitation within the ferrttic regions, and the size of the austenite-martensite islands is reduced. Higher cooling rates (of the order of 45°F/min, i.e., 25°C/min) lead, as is shown in Figure 3.4£>, to the formation of acicular ferrite which recalls martensite and in which 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 the cementite are identical to those prevailing in tempered martensite. This morphology is typical 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). X40.000 Fig. 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 MMV. (Ill,HSIRfNC.ru S TLLLS

„•) Replica - 4.5CF min (2,5=C;minK -3000 b) Thin foil — 4.5°F/min (2.53C/minl. v 10.000 Fia. 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. In this 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 contains appreciable amounts of martensite. Following slower cooling (Fig. 3.5a), precipitation of very fine carbides is observe \ within an acicular ferriie. These findings were confirmed from examination of thin foils; in particular, carbide particles oriented in one direction are 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 to structures similar to that of lower bainite. The morphology of ihe carbides changes according to whether they precipitate within martensite or bainitic ferrite; this difference is

MEAH COOLING RATE BETWEEN U70 AND S30-F."F/minuH TEST TEMPERATURE,"': nf raL io' i -120 -60 -IS 0_; 40 80 120 MEAN COOUNG RATE BETWEEN80OAND50O*C,*C/minute io' sL SL

-200 -100 - -0:•.-.. -•••-• 100--, V 200 TIME FOR COOLING BETWEEN 1470ANDgM'FleOOANDSOQ'CKminules TEST TEMPERATUBE,°F

Fig. 3.6. — Effect of cooling rate on the mechanical properties of Fig. 3.7. — Temperature dependence of Charpy V-notch impact HP 9-4-20 and HP 9-4-30. After D. COUTSOURADIS et al. [3.11]. strength of HP 9-4-20 and HP 9-4-30 after cooling within the First row of data points corresponds to martensitic range, second and bainitic range. After D. COUTSOURADIS et al. [3.1]]. third to bainitic range. Austenization : 1525=F (83O°Q - 30 min. 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 9-4-20 and HP 9-4-30 is shown in Figure 3.6. To a first approximation, and with the exception of the ultimate tensile strength of HP 9-4-30, the properties vary moderately with cooling rate. It can be seen that the continuous-cooling bainites in these steels, which, as pointed out above, are comparable to isothermal lower bainites, behave simiiariy To st-lf-ternpered martensite (cf. Section 3.5.1). As shown in Figure 3.7, which is a plot of impact strength versus 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 low temperatures.

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 about 465°F (240-C), by austenizing at 1500°F (8I5=C), hot quenching into a salt bath at 450 to 75O°F (230 to 400°C), holding for up to 16 hours, and finally oil quenching [3.14]. The bainitic reaction starts after 7.5 to 15 minutes at all transformation temperatures, and stops within two hours of isothermal holding, although this does not necessarily mean that the final structure consists of 100% bainite. For example, the end of the bainitic transformation 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 reaction temperature to 475CF (245°C) produces 98% bainite and 2% retained austenite. The structure generated by holding at 450°F (230°C), i.e. below the Ms temperature, is a more complex mixture of tempered martensite, bainite and retained austenite. The microstructural features of the bainites thus formed vary with transformation temperature. The bainite obtained between 450 and 600°F (230 and 315°C) exhibits the typical lower-bainite morphology. It should be noted that this temperature range is that within which cementite precipitates during tempering of martensite [3.15] and that, as already mentioned, there is a great structural similarity between lower bainite and tempered martensite. The bainitic plate size increases as the reaction temperature is raised. A marked transition occurs between 600 and 650°F (315 and 340°C). At 650°F and above, the bainite consists of ferrite plates with strings of cementite that tend to lie parallel 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 becomes coarser 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 presented in Figure 3.8 as a function of reaction temperature. The yield and ultimate tensile strengths decrease progressively as the reaction temperature is raised from 450 to 700°F (230 to 370°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 to 650°F (260 to 345°C), and less than 1 % in those formed at 450, 475 and 700°F (230, 245 and 37O°C). A reduction in area of 35 to 45 % and an elongation of 7 to 8 % are observed for all these reaction temperatures.

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

H 7.-1 -- energy per unit area determined on pre-eraekud substandard-thickness Charny V-nolch REACTION TEMPERMURE.'C specimens I2.ih5 • • 0.394 • 0.08 in., i.e. 55 < \O • 2 mm>. ZO ??5 30Q 325 350 A» -• critical stress intensity factor associate-! with the initiation of unstable fracture in slow bend 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 holding temperature on the mechanical properties of HP 9-4-45. After D. KAUSH et at. [3.14]. Austemzation : 1500F (815C) - 30min: bainite reaction time 6 h, followed by oil x 10.000 x 40.000 quenching, refrigeration in liquid nitrogen, and Fig. 3.9. — Thin-foil electron micrographs of HP 9-4-20 after martensitic quenching. double tempering (1 - ! h) at 400 F (2O5°C). 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 reaction temperature is raised from 450 to 600CF (230 to 315°C) (Fig. 3.8), but the transition from lower to upper bainite is accompanied by a decrease in WjA as well as a reduction in strength; 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 the critical stress intensity factor {Kv). It is thus evident that the lower bainites offer attractive combinations of strength and toughness, whereas the properties of the upper bainites are clearly disadvantageous. A reaction temperature of 500°F (260°C) gives a yield strength of 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 martensite appears to be controlled, at least partly, by the carbon content of the steel. As a matter of fact, the high-nickel steels considered here, with the possible exception of the high-carbon grades (.-.g. HP9-4-45), exhibit a microstructure of dislocation-rich lath martensite in the as-quenched condition [3.9, 3.11]. The higher Ms temperature due to the cobalt addition results in a martensite which is virtually free from austenite, but in which some self-tempering has occurred. Thin-foil electron micrographs of HP 9-4-20 after martensitic quenching are presented in Figure 3.9; it is seen that the cementite particles are arranged in a Widmanstatten stiucture typical of tempered martensite. Increasing the austenizing temperature results in a larger austenite grain, and hence in a greater martensite lath size. As a result, the yield strength decreases, as is illustrated in Figure 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 strengthening through self-tempering) [3.16]. Double austenizing treatments have been shown to give better combinations of strength and toughness because they make it possible to optimize both the austenite grain size and the amount of dissolved carbon [3.7]. In the case of

28 3. CARBIDE-STRENGTHENED STEELS — PHYSICAL METALLURGY

5Ni-Cr-Mo steels with carbon contents between 0.15 and 0.25% and cobalt contents from 0 to 8%, a double austenizing treatment of 1 hour at 1650°F (900°C) 4- 1 hour at 1600T (870uC) raises the room-temperature Charpy V-notch impact strength by 3 to 13 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 the carbon and cobalt contents [3.7], Double austenizing treatments have also been recommended for welded assemblies of the lONi-Co-Cr-Mo steels [3.8]. Prior to welding, the 1 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 the various types of steels considered in the present chapter, and their effects on mechanical properties. 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 tempering for 2 hours at 570°F (300°C), areas of acicular ferrite containing abundant cementite precipitates can be distinguished (Fig. 3.11a). Tempering at higher temperatures does not significantly modify this structure, but the triaxiality of the cementite needles becomes more marked; for tempering temperatures between 750 and 930°F (400 and 500°C), M7C3 chromium 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 made of untwinned needles; however, the extent of self-tempering is less, on account of the lower Ms point of this steel as compared with HP 9-4-20. On tempering at 570, 750 and 930°F (300, 400 and 500°C), abundant precipitation of very fine carbides, which do not appear to coalesce, occurs within the martensite (Fig. 3.lib). The following crystallo-

AUSTENIZING TEMFERATURE,°C 800 650 900 & 165 1 1 e S z x 2 5 160 \ IIOO!

• •

Q 1S5

! 150 \

MM 150O 1600 AUSTENIZING TEMPERATURE,°F

x 40,000 10.000 Fig. 3.10. — Effect of austenizing temper- a) HP 9-4-20 — Thin foil. b) HP 9-4-30 — Replica. ature on as-quenched yield strength of 10Ni-8Co-Cr-Mo steel. After T.B. Cox Fig. 3.11. — Electron micrographs of HP 9-4-20 and HP 9-4-30 after and A.H. ROSENSTEIN [3.16\. Austenizing quenching and tempering. After D. COUTSOURADIS et at. [3.11]. time 1 h, followed by water quenching. 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 TEMPERING TEMPERATURE,°C • 300 400 500 300 <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 steels A,,-, .„,„.,;„„ 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 of two experimental and two industrial 9"oNi steels quenched within the martensitic range and tempered for 2 hours. After. D. COLTSOURADIS et al. [3.11). graphic orientation relationships between cementite and ferrite in these steels have been established 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, the precipitate formed on tempering at temperatures up to 35O°F (!75°C) ;s the hexagonal z carbide; the carbide formed on tempering between 400 and 1100"F (200 and 60C°C) is cementite [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 fall when the tempering temperature is raised from 750 to 930°F (400 to 500oC); this already provides some indication of the relative insensitivity of the steel to heat treatment. This is 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 570 and 930JF (300 and 5OO'JC). It is seen, in particular,-1 that the yield strengths of the 0.3 and 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 single tempering 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 double tempered for 2 • 2 h at about IO00F (540'C). The refrigeration step for HP 9-4-30 produces an increase in yield strength of about 15,000 psi (100 MN/m-) [3.17]. The stability 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 secondary carbides.

3.5.2. 5Ni-Cr-Mo Steels As staled in the introduction, part of the developmental work on lONi-Co-Cr-Mo steels was devoted to several 5%Ni grades. This included the determination of the effect of tempering temperature on the yield strength and Charpy V-notch impact strength of 0.25C-5Ni-1.5Cr-0.5Mo steels containing increasing amounts of cobalt [3.6, 3.7]. In the as-quenched condition, the microstructure of these steels is once more typical of lower bainite or self-tempered martensite. Tempering in the 600 to 900°F (315 to 480cC) range results in the formation of FejC grain-boundary films, which are responsible for the low toughnesses measured in this range (Fig. 3.13). After tempering at I000°F (540~C), the grain-boundary Fe3C carbides are discontinuous as a result of spheroidization, hence the increase in toughness; at this tempering temperature the steels .are reported to contain both 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 the presence of a very fine carbide precipitate in the matrix. This is believed to be responsible

TEMPERING TEMPERATURE. °C 100 20D 300 400 500 700

200 400 500 800 1000 1200 1400 TEMPERING TEMPERATURE,0 F

Fig. 3.13. — Effect of cobalt content on tempering response of 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 « V STABLE — TRANSF.-J500 ON I ON COOLING COOLING

6m BOO 1D0O 1000 1200 1(00 TEMPERING TEMPERATURE, °F

Fig. 3.14. — Effect of carbon content on tempering response of 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 Ni3Mo are liable to form in the !0Ni-8Co-Cr-Mo steel during heat treatment. However, the strengthening contribution of Ni3Mo was only found to become significant on overageing, while the observations made during the initial stages merely confirmed the homogeneous precipitation of carbides generally held to be responsible for the steel's strength. The present 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 contents in relation to tempering temperature are summarized in Figure 3.14. The hardness decrease observed when the tempering temperature is raised from 400 to 800T (205 to = 425 C) is associated with the precipitation of M3C (where M = Fe, Cr, Mo) in a Widmanstatten structure. Between 800 and 1000°F (425 and 5^0°C), a secondary hardening peak is observed; this is related to the dissolution of M3C and the formation of a fine dispersion of (Mo,Cr)2C needle-like carbides which nucleate on dislocations and grow at the expense of M3C until the latter eventually disappears. From 1000 to 1200°F (540 to 650°C), where the hardness drops sharply, the M2C carbides coarsen rapidly, the substructure 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 to 705 C), the hardness starts to increase again, because the auslenite formed upon tempering now has a lower alloy content and transforms completely lo mariensite upon quenching. Finally, above I400F (760 C), complete austenization occurs and the hardness returns to 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 on dislocations, 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 tempering temperature is raised are shown in Figure 3.16 for the 0.13C grade. The fine dispersion of dislocation-nucleated M^C carbides is enhanced by the presence of cobalt, which increases the dislocation density of the matrix. Figure 3.17 illustrates this effect in the case 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 the recrystallization 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 on Fig. 3.16. — Effect of tempering Fig. 3.17. — Effect of cobalt on recovery of carbide precipitation in 10Ni-8Co-2Cr-lMo temperature on carbide distribution dislocation substructure in tempered 0.12C- steels tempered for 1 hour at 950°F in 0.13C-10Ni-8Co-2Cr-lMo steel. 10Ni-2Cr-lMo. After G.R. SPEICH et al. [3.9). JSKTC). After G.R. SPEICH et al. [3.9]. After G.R. SPEICH et al. [3.9]. Condition : I55OCF (845°C) - 1 h. irAustenization : 155O°F (845°C) - 1 h, W.Q. Austenization : 155O°F (845°C) - 1 h, W.Q. W.Q. +950°F(5103C)-12h.A.C.

33 1.HUM V i ONI \lM\li HR.H-SIKI NlHH STIflS

TEMPERING TEMPERATURE.'C 100 200 300 «X) 500 600 700

1300

ai2C-ONi-6C

noo :

VWOg 3 SCO Z !0.12C-l0Ni-8Co-2Cr

1400 1000 1200 800 TEMPERING TEMPERATURE.'F ffl2C-10Ni-aCo IK. ?.|S. HTeet of tempering temperature on austenite formation m H)"f,Ni -tceK. -\Her Ci.K. SPHIH el ill. [3-'-*\. Q12C-10N1 •Vii-tcni/ation : 1550 K (S45 O - I h. W.Q. 600 Tempering time : 1 hour.

500 DNi : hiu 3.11). I tl'ect of individual alloying elements on tempering response of 10",,Ni sieelv After OR. SPHI ii ft at. \3.t\. RT " 200 «X) 600 600 1000 1200 VO0 AitMeni/ai-.nn : 1550 V (H45 Cl - ! h. W.I). TEMPERING TEMPERATURE.'F Tempering time ; 1 hour.

3.5 4. Retained .-Ui.slcnite Formation of austcnite during tempering and its tetention upon quenching from the tempering temperature have been studied in the case of a 0.12C-iONi-8Co-2Cr-!Mo and a O.I2C-IONi steel [_?.<>]. Figure 3.18 shows the amount of austenite retained at room temperature after tempering for 1 hour. No retained austenite is observed in either steel after full austenization and quenching: ii only starts to occur at temperatures exceeding the secondary hardening peak, in amounts which increase as the tempering temperature is increased to 1200 F (650 C) and then decrease to zero at 1300=F (705"C). Prolonged exposure (up to 20 h) at 950 F (510 C) does not result in any austenite Eormation, whereas exposure 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 a low-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 yield strength 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 is observed 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 strong influence on this property after tempering. Cobalt and, to an even larger extent, chromium increase the steel's resistance to softening, whereas molybdenum contributes to secondary hardening. 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

160, "150 160 170 180 190 200 0.25 0.30 0.35 0.40 0.30 0.35 0.40 0.27. OFFSET YIELD STRENGTH, CARBON, wt.% CARSON, wt.V.

Fig. 3.20. — EITect of carbon content on notch toughness - yield strength Fig. 3.21. — Effect of carbon content on mechanical relationship for IONi-8Co-2Cr-lMo properties of HP 9-4-X steels in tempered martensitic steel. After G.R. SPEICH el at. [3.9]. condition. After J.S. PASCOVER and S.J. MATAS [3.2]. - \ vniiation : 1550'F (S45C) - I h, W.Q Condition : austenization, oil quenching, refrigeration, Innpering for indicated times at 950"F (510 C). 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 carbon contents was evaluated from the start, since this element is obviously responsible for basic hardening 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 shows the toughness vs. yield strength relationship for 10Ni-8Co-Cr-Mo steels with different carbon 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 occurs at this temperature and of the simultaneous increase of both yield strength and toughness that results under some tempering conditions. The effect of carbon on the strength and toughness of the HP 9-4-X steels is illustrated in Figure 3.21. The strength increases with increasing carbon content, whereas the toughness decreases; this effect is less marked for the higher-carbon grades. Increasing the carbon content aiso decreases the ratio of notched tensile strength to yield strength, irrespective of the strength level [3.2]. Since the general effect of carbon is thus to increase strength but to decrease toughness, it is evident that the optimum carbon content must be chosen in 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 hardenability. Another, and even more important one, is to lower the transition temperature so that 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 strength itself, is probably associated with the fact that nickel enhances cross-slip at high strain rates and/or low temperatures [3.19]. The effect of nickel in lowering the M? temperature and thus in increasing the amount of retained austenite is counterbalanced by means of cobalt additions.

35 COBALT-CONTAINING HIGH-STRENGTH STEELS

fig. 3.22. — Effect of silicon on toughness of TEMPERING TEMPERATURE/C HP 9-4-45 in tempered martensitic condition. 200 300 400 500 1.00 r After J.S. P*si:ovm and S.J. MATAS J.?..'l. Condition : uustcnizatinn, oil quenching, refrigeration, and 2 -• 2 h tempering at indicated 0.751 temperatures.

0.25

300 400 500 600 700 800 900 ITOO TEMPERING 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, special melting 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 characteristics is concerned. However, its presence in substantial amounts does not contribute to toughness as nickel does indirectly. The upper limit for HP 9-4-25 is 0.35 %Mn, and most production 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 steels is 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 of the latter as compared with the former at all tempering temperatures are evident, as is the decrease 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 that the chromium and molybdenum contents be kept to a minimum in HP 9-4-45 when the highest level of toughness was required. These two elements do not have a detrimental effect 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-temperature properties. The role of chromium and molybdenum has been investigated in greater detail in the case of the IONi-8Co :,teel system [3.9]. Chromium was shown to shift the secondary hardening peak 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 hardening is 2%; higher levels result in too rapid coarsening of the M2C alloy carbides, while lower contents are associated with problems of retained austenite, due to a higher secondary, hardening temperature. Molybdenum at the 1 % level is responsible for the occurrence of a marked secondary hardening peak in the 0.12C-10Ni-8Co-2Cr steel; increasing the molybdenum 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 rates for both the pearlite and bainite transformations; however, its major function is to act as a grain 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 mentioned in preceding sections. In brief, it ircreases the Mg temperature, refines the martensitic structure, and leads to retention cf the dislocation substructure at higher tempering temperatures, giving a finer precip lation of dislocation-nucleated alloy carbides. The incidence of these structural features on the tempering response of cobalt-containing steels will now be discussed in greater detail. The effect of increasing cobalt contents on the yield and impact strengths of a 5Ni-Cr-Mo steel containing 0.25 %C was shown in Figure 3.13. Addition of 4%Co enhances the secondary hardening effect, so that the yield strength is raised on tempering in the 800 to 1000 F (425 to 540C) range [3.6]: increasing the cobalt content to 8",, increases the yield strength at all tempering temperatures. On the other hand, the notch toughness is lowered, especially in the region of secondary hardening, so that the strength/toughness relationship for the Co-containing teels tempered at I000=F (540cC) is poorer than that obtained for the Co-free one tempered at 400°F (205°C) [3.6, 3.7]. In the case of the former, 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/toughness combination.

The effect of up to 8% cobalt on the tempering behaviour of a carbon-free, 10°oNi steel is to increase the hardness through a small solid-solution effect which appears to be retained over the whole tempering range [3.9]. Both the Co-free and the Co-containing steels 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 a 0.12C-10Ni steel increase the hardness and strength values at almost all tempering temperatures. In the presence of carbide-forming elements (Fig. 3.23), cobalt additions

TEMPERING TEMPERATURE,*C . WO 200 300 - 500 600 160 5-1

2QQO

180 £

160 <

U0E

120?

100?

0 SO 3 40 3. As quenched 200: UIO 600 ._ BOO 1000 1200 1EMPERIN6 TEMPERATURE.°F

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 OS I \l\l\ii MliiH M KIMJ III SlLLUs

result m much higher attempered yield strength values; in particular, a sharper secondary hardemns: 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 interpreted in 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 be achieved m the cohalt-modiiied lONi-Cr-Mo than in the eohah-moditied 5\i-Cr-Mo steels referred to earlier on: on lhe one hand, the increase in yield strength at the secondary hardening peak is much more significant in I he former steels than in the latter; on the other, the toughness trough for the IONi-Co-Cr-Mo steels occurs ai a lower tempering temperature than does the yield strength peak, with the result thai, at the temperature corresponding 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 condition can be considered as being derived from three primary strengthening mechanisms (see Chapie.' 2. Sections 2.3.6 and 2.-J.1) : (1) substruetural strengthening, resulting from the high dislocation density of the martensite. (2l carbide precipitation hardening, and I3I suh>tiiL,.ional-element solid-solution strengthening. In steels that are subject to secondarv hardening, it is important to retain the dislocation substructure at temperatures as high as possible since, in addition to its inherent strengthening effect, this structure will result in a liner dispersion of dislocation-nucleated particles and hence in increased precipuat.on hardening. An estimation of ihe magnitude of the different strengthening mechanisms operative in a 0.12C-10Ni-KCo-2Cr-IMo steel has shown that, for a yield strength of IS 5.000 psi (1200 MNm-'l. carbide strengthening contributes 45",1, sub- structural strengthening 37",,. and solid-solution strenthening 16";,, the balance being accounted 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 various elements in determining the strength of tempered IO",,Ni martensites. Although alloying is intended mainly to control the steel's tempering behaviour, the high nickel and cobalt contents give rise to some solid-solution strengthening. Carbon obviously increases the strength of the steel over the entire tempering range through an increase in the amount of carbides and the consequent decrease in intercarbide spacing. Chromium reduces ihe rate of softening but does not give a secondary-hardening peak in the absence of molybdenum. The latter element (in association with chromium) is responsible for secondary hardening 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 be attributed to two effects, vhich probably occur simultaneously ; (I) cobalt increases the

Fig. 3.24. — Interface mod>l showing the effect of cobalt partitioning Y on carbide growth rate in steels. After V.K.. CHANDKOK et al. [3.22]. "DIFFUSION PATH FROM INTERFACE-^ Cci, is the concentration profile for cobalt; or and ac(Co) are res oj partitioning b) no partitioning pectively the activities of carbon in the absence and presence of cobalt.

38 i. CARBlbil-STRENGTHENED STEELS — PHYSICAL METALLURGY

dislocation density and retards the recovery of the dislocation substructure, so that the number 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 a liner dispersion of, the carbide precipitates, as shown particularly clearly in the case of high-speed steels [3.22, 3.23]. The elfeet of cobalt on the activity of carbon should also lead to a reduction in the growth rale of the carbides, in conjunction with the partitioning of 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 activity gradient at she interface (hence the decreased carbide growth rate) [.''.22]. Thus, cobalt enhances both substructure and carbide precipitation strengthening, i.e.. the two factors which have just been shown to account for the largest part of the strength in the case of the 10Ni-8Co-Cr-Mo steel. However, the mechanisms mentioned in the preceding paragraph as controlling the nucieaiion and growth of carbides do not rule out a possible effect of cobalt on the diffusion of carbide-forming elements [3.24] or on the surface energy of the carbide,ferrite interface [3.22]. The increased activity of carbon due to the presence of cobalt should also be responsible For the easier self-tempering of cobalt-containing carbon martensites, and possibly for the increase in Mf temperature in carbon-depleted zones [3.22]. Toughness. As mentioned earlier, the role of nickel with respect to toughness is mainly to decrease the transition temperature, and hence to ensure retention of high room- temperature toughness at high strength levels; there is. however, a limitation to the use of nickel as an alloying element, which is primarily related to its tendency to increase austenite retention or reversion. The effect of the other alloying elements on toughness can be interpreted in terms of their role in promoting the substitution of the much tiner MiC for MjC: ductile fracture occurs by nucleation of voids at particles, and such nucleation 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 simultaneously high 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 essentially along two lines : (1) the development of low-carbon martensites strengthened by intermetallic compounds, and (2) the development of steels with the highest carbon contents compatible with high toughness and good weldability. The present chapter has been concerned with the second type of development which characteristically involves the use of high nickel additions to achieve high toughness. Imperatives regarding the steels" transformation structures {e.g. restriction of the retained austenite content) led to associating this element with cobalt: the levels which were selected on the basis of property requirements, and possibly cost, were 4 and 8"oCo for the HP 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-forming elements chromium and molybdenum. In the lONi-Co-Cr-Mo steel, the optimum chromium and molybdenum contents were determined as 2 and 1 %. respectively, whereas in the HP 9-4-X steels, they were held to 1 % or less depending on the carbon content. The latter 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 rather low (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 greater than wH'i-llv. j.t.. MU) at \lcki strength le\els up to 21)0,000 psi (1400 MN..-m:). The carbon level of HPlM-\ steeK varies from 0.20 to 0.45",, depending on the combination of toughness, weldability and strength that is required. The aood properties of these steels in the tempered marten-.itc condition result from the tact that their >.iruciure> are strengthened b\ a line dispersion of alloy carbides which substitute at least partially for the embrittling cementite precipitate. This process is associated uub a relali\e inseii>itivitv of pri.perties to tempering temperature in ihe case of the HIJlJ~i-\ steels and. in that of the !U\i-(.'o-C"r-Mo steels, with the existence of a secondary ha; Jening range uver which both the yield strength and toughness increase. The former sleek, particularly vhe HP"-M-45 grade, also exhibit excellent combinations of 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-X steels. Data on the lONi-SCo-Cr-Mo steel will also be included, although this steel is not so well documented, due partly to the restricted distribution of some of the reports on iti development. However, in view of the similarity between the IONi-8Co-Cr-Mo steel and the 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 absence of specific data in many cases.

4.1. Primary Processing The HP 9-4-X steels are produced by electric-furnace air melting pins consumable-electrode vacuum-arc remelling. The optimum combinations of properties are obtained using the VAR-CDOX process (vacuum-arc remeiting plus carbon deoxidation): other melting practices (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 excess carbon, followed by consumable-electrode vacuum-arc remeiting using a cold-walled crucible [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 shown that steels prepared in this way tend to be cleaner than those that are air melted, or air melted plus vacuum remclted. and to have lower gas contents. Both factors lead to an improvement 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 carbon deoxidalion [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 the absence of significant amounts of aluminium: in Si/AI deoxidized material, grain coarsening has not been observed below 2150 F (1175 C) [4.6]. When forging VAR-CDOX ingots of HP 9-4-45 and HP 9-4-25, a maximum reduction of 75% and a maximum finish temperature 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 regards hardenability, essentially no drop in hardness with distance from the quenched end of Jominy bars has been observed for either the 0.25C or the 0.45C grade; this indicates

40 4. CARBIDE-STRENGTHENED STEELS — PROCESSING AND PROPERTIES

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 at Temper 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 al i 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 hardenability of HP 9-4-20 and HP 9-4-30 appears to be even larger : according to the CCT curves given in Chapter 3 (see Section 3.2.1), the rounds that can be fully hardened by water quenching 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 martensitic treatment (the A/s temperature is about 450°F, i.e., 230°C), but this must include a refrigeration step to minimize the retained austenite content. It can also be subjected to a bainitic treatment, which produces higher toughness but a somewhat lower yield strength. The bainitic transformation in the 450-650T (230-345°C) temperature range is completed in less than 2 hours; the 4 to 8 hours mentioned in Table 4.1 are primarily a precautionary measure, 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 the benefits 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 high transformation temperatures allow self-tempering of the mariensite to occur, leading to good toughness in as-quenched or as-welded material. It is not possible to produce a fully bainilic structure in this grade over a wide range of reaction temperatures as was the case for the higher-carbon steels; accordingly, the only heat treatments used are those that will produce a martensitic structure. However, such treatments are adequate since the toughness of the tempered martensite is sufficiently high and would not be significantly improved 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 tempered condition. The standard heat treatment listed in Table 4.1 involves a double austenizing treatment [4M], though this has occasionally been replaced by a single treatment at 1550 or 150H F tS4? or i>15 C) [4.11, 4.12}. Hardenabilitv is more than adequate for plate thicknesses 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 obtained in the HP 9-4-45 and HP 9-4-30 grades and are associated with good toughnesses, particularly 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 tensile strength level of 270.000 psi (1860 MN/m-) [4./]. The lower-carbon grades are conspicuous foi 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 typical A',,, 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 combi nation of strength and fracture toughness than does HP 9-4-20 [4.9]. K]c values determined according to ASTM specifications are not available for this steel, because of the large specimen thickness required. Estimations based on results of tests on sub-thickness specimens are in excess of 250.000 psiy in. (275 MNm-*'-) [4.13], and Ku. values as high as 300.000 psi\ in. (330 MNm ••<•'-) have been cited [4.5]. The effect of melting and clenx'ulaiiun practice on tbe strength/toughness relationship has been investigated in the case of the HP 9-4-45 steel. Figure 4.1 clearly shows the improved toughness of vacuum carbon-deoxidized heats as compared with air-melted Si/Al deoxidized ones. As regards the effect of heat treatment. Figure 4.2 shows the influence of tempering temperature 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

U T.S.. 0.2' „ Y.S.. Elong. R. C.V.N impact, Kir, Hard- Material ness, Ref. and condition * in 2 in.. in A.. 10-' psi MS m~ HP psi MWm* ft.lb J SO^psi v in- ^ Re

HP9-4-45 (muncnsitic) 2X0-300 1930-2070 245-260 16F.O-I790 6-10 20-35 10-12 14-16 50-70 55-77 51 [4.7] HP9-4-45 (bainuicl 260-280 1790- IV30 220-235 1520-1620 12-14 40-50 20-30 27-41 65-95 71-104 51 [4.7] HP9-4-3O(martsnsilic) 220-240 1520-1660 190-200 1310-1380 12-16 50-60 20-25 27-34 100-120 110-132 44 [4.7] HP9-4-30 (bainuicj 220-240 1520-1660 190-200 1310-1380 12-16 50-60 25-30 34-41 120-135 132-141 44 [4.1,4-7] HP 9-4-25 195-210 1345-1450 178-192 1225-1325 15-Io 55-65 32-40 43-54 > 140 > 154 42 [4.7] HP 9-4-20 190-215 1310-1480 180-195 1240-1345 14-20 61-71 51-72 68-98 155-175 170-192 42 [4.7,4.8] lONi-SCo-Cr-Mo 196 1350 185 1275 22 71 81 110 300** 330** 45 • [4.5]

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

42 4. CARBIDE-STRENGTHENED STEELS — PROCESSING AND PROPERTIES

YIELD 5TR6NGTH: MN/m> TEMPERING TEMPERATURES I5oo; •- •••- -is - lam

910 . 220 r -230 , • ;2A0 YIELD STRENGTH, tfpsr:' As « 400 800 1200 » 400 800 1200 quenched TEMPERING TEMPERATURE.T

Kig. 4.1. — EITecl of deoxidation practice on the strength/toughness relationship in HP 9-4-45 sheet, a) HP 9-4-20. Single 4-hour temper. l» HP 9-4-30 (martensitic). l-atigue cracked centre-notched specimens 0.180 in. Double 2^2 hour temper. !4.h mm) thick were used in determining the longit- Fig. 4.2. — Effect of tempering temperature on properties of HP 9-4-20 and udinal notched strength. After SJ. MATAS [4.3]. HP 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 bainitic conditions 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 the yield 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 thermomechanical treatment has been investigated for the HP 9-4-45 (martensitic) and HP 9-4-25 grades [4.17]; the effect of strain tempering was evaluated in both cases in terms of the extent of cold work (roiling) and the tempering temperatures prior to and after cold

YIELD •STRENGTH, MN/m' 1000 1200 1400 1600

Fig. 4.3. — Relation between strength and impact toughness of 0.5 in. HP 9- 4-45 plate heat-treated to bainitic and tempered martensitic structures. After J.S. PASCOVER and S.J. MATAS [4.16]. Data points from left to right corres- KO 160 180 MO 220 240 260 pond to decreasing bainitic trans- 3 formation temperatures or decreasing YIELD STRENGTH, B psi martensite tempering temperatures.

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

PRETEMPERING TEMPERATURE.'C RETEMPER1NG TEMPERATURE/C

....200'.. X» -400^. 500 300 j}- 400: ;:500 6CXT

0 n 20 30 40 300 500 700 900 1100 'RT 200 " «0 600 . SOO WX) 1200 THICHNESS REDUCTION, V. . PRETEMPERING TEMPERATURE. "F RETEMPERING TEMPERATURE,-F .

EtTccl i)l' .mioum of deformation Ipre- »*>) Effect of pretcmperinK ;cmpc: .i.'-irc i-5'\, ri Effect of retenperiny tcmpemturL' (pre- d posl-U'mrcrmg ai 4\XT F. i.e. 205 Ci. deformation, no posl-deformult'Jit ijinpcring}. 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 was found 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 low temperature (around 300 F. i.e.. 150 Cl resulted in much more pronounced strain hardening than did pretempering at high temperature (about 1100'F, i.e., 595QC); this can be 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 in yield strength, independently of the degree of deformation and the tempering temperature over the 200 to 400 F (95 to 2O5X) range. All these effects are strikingly summarized in Figure 4.5a which shows the KUJY.S. relationships for the steel both after the conventional martensitic heat treatment and after subsequent strain tempering; although the two curves initially 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 acceptable Fracture toughness (around 50,000 psi\ in., i.e. 55 MNm"3'2). HP 9-4-25 can also be strain tempered to extremely high yield strength levels, and these prevail in conjunction with higher toughnesses than are obtainable with HP 9-4-45 (Fig. 4.5b).

YIELD STRENGTH MN/m* _2«0 Z300 1S00

42D 180 220 YIELD STRENGTH, 103psi

Fig. 4.S. — Strength/toughness relationship for HP 9-4-45 (martensitic) and HP 9-4-25 in both the conventionally heat-treated and strain-tempered conditions. After D. KALISH el al. [4.17]. 4. CARBIDE-STRENGTHENED STEELS - PROCESSING AND PROPERTIES

YIELD STRENGTH MN/m* 1200 1400 1600 1600 2000 2200 260 T 10Ni-8Co-Cr-Mo 240

250 220

200

ISO 200

. 160 ID 20 30 40 50 THICKNESS REDUCTION, % ^•140 . 18Ni-Co-Mo sa MARAG1NG STEELS ; I50 : Fig. 4.6. — Effect of strain tempering on mechanical properties of HP 9-4-45 in the bainitie 120 condition. After D. KALISH el al. [4.18]. Transformation to bainite at 500'F CttfTC); pretempering at 400°F (2O5=C); 100 no post-deformation tempering. 100 80

50 40

Fig J.7. — Comparison of strength/toughness relationships for the HP9-4-X and other high-strength steels. 20 ISO 180 ZOO 220 240 260 280 300 Alter [4.8] for HP 9-4-20, J.S. PASCOVER [4.15] for HP 3 9-4-30, J.H. GROSS [4.13] for I0Ni-8Co-Cr-Mo, and J.S. YIELD STRENGTH. 10 psi PASCOVER and S.J. MATAS [4-16] for the other steels.

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 bainite leads to increased strength and decreased toughness (Fig. 4.6). Ausforming HP 9-4-45 has also been shown to increase the strength properties and to decrease toughness {4.17}. However, due to the fibering effect associated with rolling in the austenitic condition, anisotropy in fracture toughness develops; the resulting longitudinal strength/toughness curve is close to that for the steel in the conventional and strain-tempered martensitic conditions, but the transverse strength/toughness data points lie below it. This section can be usefully concluded by comparing the strength/toughness relationships for the HP 9-4-X steels with those of other high-strength steels (Fig. 4.7). It is apparent t!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 2 3 2 (1310 MN/m ), HP 9-4-20 exhibits Klc values of up to 180,000 psi\/in. (198 MNm" ' ), while HP 9-4-30 appears to be at least as good as the 18%Ni maraging steels at strength 2 levels of "200,000-210,000 psi (1380-1450 MN/m ). Although comparable K]c data are not available for the 10Ni-8Co-Cr-Mo steel, the values quoted in Table 4.2 indicate that, at the 190,000 psi (1310 MN/m2) level, the curve for this steel probably lies well above the highest values for HP 9-4-20.

45 lOHM.r-tONTAlMMI HK.H-S I Rl-NCH H STEKLS

TE5T- TEMPERATURE.'C TEST TEMPERAHJRE,"C TEST TEMPERATURE,*C ,0_ 200 400 600 800 BOO .-200 " -SO • -100 -50- 0 -200 -100 0

BO r '• 60 j Q / 40 No' 20 =—-cSS!!gil>' n 0 400 800 E00 -BOO 2000 ..-300 -200 -100.. 0 BO - -»a -2oo -wo o TEST TEMPERATURE. F tEST 1EMPERA1URE/F TEST TEMPERATURE, T

Fig. 4.8. — Temperature dependence ii) Tensile properties and frac- h) Impact strength of HP 9-4-45 and HP 9-4-20. Respectively of tensile properties of HP 9-4-20. ture toughness of HP 9-4-20 after J.S. PASCOVER and S.J. MATAS [4.16]. and R.T. AULT [4.14], plate. After C. VISHMEVSKY HPy-4-45 was treated to 250.000 psi (1720 MN/m') After R.T. ALLT [4.14]. ..id E.A. SrEictRWALO y.iV]. U.T.S. in both the bainilie and martensitic conditions: Condition: P00F (93OC)-lh. A.C.- 1500 F Condition : standard ht. treat. HP 9-4-20 received the conveniional heat treatment. ; <8I5C) - th. W.Q.- IO5OF (565 C) - 2h. A.C. 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 is rather scarce. Figure 4.8 illustrates the effect of test temperature on the mechanical properties of HP 9-4-20 after its standard heat treatment. Regarding the properties of '>ese steels at cryogenic temperatures, F'gure 4.9a shows those determined on HP 9-4-20 plate in the heat-treated condition. The ultimate tensile strength is seen to increase from 200,000 psi (1380 MN/'m2) at room temperature to

-20C 10* TO5 10s 1000 CYCLES TO FAILUHE

Fig. 4.11. — Smooth and notched fatigue behaviour Fig. 4.10. — Stress-rupture and 0.2",, plas- of HP 9-4-20 (Moore rotating-beam fatigue tests on transverse specimens). After R.T. AULT [4.14]. tic deformation creep curves for forged 1V S HP 9-4-20 bar (transverse). After [4.20]. Specimens tempered at IOOO F (54O C) lo ultimate tensile and yield strength levels of 210.500 and 187.000 psi f 1450 and 1290 MN/m*), respectively.

46 4 < \RHIDI -SIKI-N(,!III-N[.D Ml IIS PROCESSING AND PKOPF.R7 li:s

Fig. 4.12. — Dependence of fatigue crack growth rale in air on stress intens- STRESS INTENSITY FACTORRANGE^K.MNmW ity factor range lor HP 9-4-20 and jjj>.'";--''•'••'•' 50 '1-,-X' '100 :-•-...• 200 .A>-i'.i • lONi-XCo-Cr-Mo steels. After [-4M]. o HI'9-4-20 ; lONi-KCo-Cr-Mii

£ 20 20 50 100 200 STRESS INTENSITY FACTOR RANGE, &K.103 psi \fin.

255,000 psi (I760MN/m-) at —32OCF (—I96°C), at the expense of a sharp decrease in fracture toughness and a slight decrease in ductility [4.19]. Similarly, the ultimate teusile 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 of temperature from —300 to 120F (—185 to 50;C) on the impact strength of HP 9-4-45 and HP 9-4-20 is shown in Figure 4.9b. As was to be expected from the typical behaviour of high-nickel steels, the decrease in impact strength with decreasing temperaiures is not characterized by a transition temperature. Finally, creep and stress-rupture data determined on forged HP 9-4-20 bars are presented in Figure 4.10. 4.2.3. Fatigue Behaviour Results of room-temperature axial fatigue tests on both smooth and notched specimens of HP 9-4-20 are plotted in Figure 4.11. Further data on the fatigue behaviour of HP 9-4-X steels will be found in the literature [4.5, 4.10, 4.20]. The fatigue crack propagation rate of several high-strength steels tested at room temperature in an air environment was shown to obey the following relationship : do/dW = 0.66 x IO~8 (A/C)2-25 [4.21]. Figure 4.12 illustrates this relationship in the case of HP 9-4-20 and 10Ni-8Co-Cr-Mo, though in 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-X steels determined on notched specimens is generally improved by factors that tend to increase its fracture toughness, such as reduced strength level and carbon content, as well as by special melting practices which minimize non-metallic inclusions [4.22].

Threshold intensity factors KUcc of several high-strength steels in 3.5 % sodium chloride solution are given in Table 4.3. It is apparent from these results that HP 9-4-45 has only marginal resistance to stress-corrosion cracking in the martensitic condition; it is definitely more resistant in the bainitic condition, though still greatly inferior to HP 9-4-30. These observations 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]

1 1

Fit. 4.13. — K\,rr stress-corrosion behaviour of HP'MOO in artificia' sea-water (3.5",,NaCl). After y.S\. 1001 0 3To 400 600 800 BOO 1200 K00 1600 1800 r> Air tesis TIME 10 FAILURE, hours • Sea-water tests • > No failure after six months I I reflect the latter's superior stress-corrosion resistance. Results of stress-corrosion tests on 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]; this also applies to the HP 9-4-30 grade treated to martensite, but not to HP 9-4-45, which was fabricated in the annealed condition. Possible operations include rolling, drawing, bending, roll forming, shear spinning, explosive forming, etc. The effect of cold working on the mechanical 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 heat treatment, their machinability is comparable to that of AISI 4340 at similar hardness levels [4.7]. Finally, the HP 9-4-X steels, in particular the lower-carbon grades, have been shown to possess excellent weldability. Data have been reported for gas tungsten arc (TIG), hot-wire TIG, 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 filler metals which closely match the base metal composition [4.26]; tbey require ho post-weld heat treatment in view of the good toughness of the self-tempered martensite and/or

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

U.T.S., y.s.. f Steel Reference 1 3 3 10^ p. i MNjiri 10 psi i 10 psiv'in. 103psi\/in. M

300 M 283 1951 236 1627 76 S3 13 14 [4.23] HP9-4-45 (martensiticl 276 1903 236 1627 69 76 15 16 » HP 9-4-45 (bainitie) 266 1834 220 1517 89 98 20(est.) 22 » 4330 V 239 1648 196 1351 103 113 25 27 » H 11 219 1510 188 1296 54 59 30 33 » HP 9-4-30 (martensitic) 231 1593 200 1379 116 127 45 (est.) 49 » 18Ni(250) maraging 259 1786 249 1717 92 101 45 49 » HP 9-4-20 205 1412 190 1310 175 192 110 121 [4.8] lONi-SCo-Cr-Mo (plate) 196 1350 185 1275 .100 330 180 198 [4.5] l0Ni-8Co-Cr-Mo (hoi pressed) 203 1400 191 1317 > 234 >257 210 231 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, joint efficiencies of 95 to 100",, have been reported [4.1). The properties of HP 9-4-20 we-ds arc equivalent to those of the heat-treated base plate [4.14]. In order to overcome detrimental effects of stressrelieving treatments on the toughness of the welds, studies on the use of modified filler-metal compositions were recently carried out; they resulted in the development, in the case of HP 9-4-20,of a completely stress-relievable system, i.e., one in which the base metal, heat-affected zone and weld are totally free from stress-relief embrittlement [4.28]. As regards the HP 9-4-45 grade, it was shown that, like other medium-carbon ultrahigh-strength steels, it must be welded in the annealed or normalized condition and then fully heal treated to develop the desired properties; using filler metals of matching composition except for a slightly lower carbon content, joint efficiencies of nearly 100",, on the basis of yield and ultimate tensile strengths were achieved, with fracture 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 comparable to those oi' the lower-carbon HP 9-4-X grades. In addition, the steel has been successfully welded by the TIG process using weld wire of composition nearly matching that of the base metal [4.9]. A comparison of weldments made v/ith iuNi-8Co-Cr-Mo and HP 9-4-20 wires showed that the most desirable weld properties are obtained using the matching weid wire and a tempering cycle. Under these conditions, tensile properties, impact strength, precracked impact energy and " apparent " fracture toughness of the weld centre and heat-affected zone are close to those of the base metal. It is stated, however, that the improvement gained through retempering is probably not sufficient to justify the additional cost 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 during fabrication or which do not preclude a final heat treatment; this is particularly so for small parts which are fabricatedby forging, machining, and possibly welding prior to heat treatment. For instance, the remarkable combination of strength, toughness and fatigue 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 spring wire 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 homogeneous armour plate because of their good hardness/toughness relationship. On account of its high structural stability up to at least 800°F (425°C), associated with satisfactory retention of mechanical properties, HP 9-4-30 has also been selected for several forged structural components on the Boeing 747 aircraft [4.25). HP 9-4-20 has also found applications as forged 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 that they can be welded in the heat-treated condition and require no post-weld treatment. Moreover, their deep-hardening characteristics, together with their high strength, high crack-propagation resistance (ATlc. and ^ijCC) and superior toughness, make them ideal for use in heavy sections in aircraft, aerospace and hydrospace applications. In fact, both 10Ni-8Co-Cr-Mo and HP 9-4-20 have interesting potentialities for the manufacture of internally or externally pressurized vessels, such as reactors and submersibles. In this respect, a great deal of evaluation work has been devoted to HP 9-4-20 by the U.S. Atomic Energy Commission and the American Society of Mechanical Engineers in connection with nuclear reactors and pressure vessels, respectively.

49 1-iOM AIMMI HICHSI RFNCilH STEMS

? Ni-Co-Mo MARAGING STEELS — PHYSICAL METALLURGY

5.1 Background The maraging steel:- are a relatively new class of ultra-high-strength sieels that dcrise their strength from hardening mechanisms other than the classical ones associated with carbon martensite. bainite or precipitation of carbides during tempering. These steels. which possess combinations of strength and toughness that are among the highest attainable in commercial alloys, are characterized by their very low carbon content and ihe use of suhsiituiionul elements to produce age-hardening in Fe-Ni mariensites. The metallurgical principle on which manging steels are based was established as early as 1939 [5.1] : the thermal hysteresis between martensite formation in Fe-Ni alloys on cooling and its reversion to austenite on heating, and ihe increase in this hysteresis with increasing Ni content (Fig. 5,1). It can be seen that the reversion temperature of a 20"v,Ni alloy is about 1IO5;F (595 C). which is sufficient to allow ageing of the martensitic matrix at about 900 F (480C). The term " maraging " was coined to indicate that the precipitation reactions that are responsible for the ultra-high strength of these steels occur on 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, the so-called 20",, and 25"0Ni steels [5.4). The A/., temperature of the alloys was controlled by 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, were precipitation hardened during ageing between 800 and 950F (425 and 510 C). Hardnesses as high as Rc 67 were obtained, and good combinations of strength and ductility at hardness levels of Rc 53-56 were reported. However, the 20 and 25",,Ni grades were soon abandoned on account of their brittleness at 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 40 60 12001 (2501 13001 13501 i«BI 15001 wt.%Coxwt.%Mo MARAGING STEEL GRADE

Fig. 5.2. — Effect of cobalt • molybdenum Fig. 5.3. — Ultimate tensile strength, product on maximum hardness of Fe-18.5 to fracture toughness and cobalt and 20.1 '•„ Ni alloys. After R.K. DK KF.R el at. [5.6). titanium contents vs. maraging steel o annealed I h al 1600 T (870"Cl, A.C. grade. After A. MAGNEE er al. [5.7]. * maraged 3 to 10 li at 800-900 F (425-4hl0C). martensites that occurs when combined additions of cobalt and moKlxlenum are made was reported in 1960 [J.6]. Figure 5.2 illustrates the synergistic effect of coball with molybdenum on the age-hardenability of Fe-18"0Ni alloys. This binary composition was chosen as the alloy base since higher nickel contents led to retained austeniie. A number of l8Ni-Co-Mo alloys were developed in rapid succession; later on. other Ni-Co-Mo compositions with less nickel were introduced. Basically, the heat treatment of maraging steels consists in solution annealing for I hour at 1500°F (815"C), although other annealing temperatures or multiple annealing treatments are used in certain cases. Upon cooling in air to room temperature, the alloys transform completely to martensite. Because of their high nickel content and the virtual absence of 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 the treatment are very small, so that in many cases the parts can be completely finished before hardening. Maraging steels can be welded without preheat in both the annealed and fully heat-treated conditions. These alloys possess resistances to hydrogen embrittlemenl and stress-corrosion cracking 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 as constructional alloys, particularly in extreme-duty applications requiring high strength- to-weight ratios, is their excellent toughness at high strength levels. In particular, they exhibit 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 typified by-AlSl 4340. The plane-strain fracture toughness (Klc) of the 18",,Ni maraging Mcels, at yield strength levels of 240,000 to 280,000 psi (1650 to 1850 MN/'ra-), is more than twice that of the best 4340-type steels. The ultimate tensile strength and fracture toughness of various maraging steel grades are shown in Figure 5.3. (Ill \l I MINI UNIMi Hit ill NIRt \t,lll

1 ollowins: the development of niaragini; steels, considerable effort has been devoted to Miuhtne 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 in Reference- -^ v and -"" •' Although the various phase transformations and age-hardening reactions 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\ have been included in the following pages.

- ! ! K,::t '•! -I/.•'('I/(If l-ii'-'hWS 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 energy of the interaction of dislocations with interstitial atom- [.vV]. It promotes stress relaxation and. as a re>ult. reduces the susceptibility of the steel to brittle failure. The marten-mc structure creates favourable conditions for uniform nucleation and distribution of mtermetailic phases during ageing, thus ensuring higher plasticity and Jiictilitv. Detailed -Indies of Fe-Ni martensites containing addition dements such as Vi. Be. A I. Mo. Mn. Nb. Zr. \V and Cu. have shown that age-hardening occurs in the t'dii in 12(10 |- i?50 ni (,50 Ci range. However, the degree of hardening may vary appreciably with changes in the nickel content [5./f>]. or with variations in the ageing or reversion reacnon kinetic- Moreover, rather strong interactions can occur between specific combinations of elements, such as cobalt and molybdenum. Nevertheless, the alloying .•lemenis can be classiiied qualitatively as "strong" (Be. Til. "moderate" 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 steels through 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 lie up residual carbon. During solidification of the steels, this element tends to segregate or to 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]. Precipitation of compounds of titanium with carbon, nitrogen and oxygen was also detected on ageing maraging steels between 700 and 900 F (370 and 480 C) [5.12\. The critical titanium content above which ductility is lowered h.ith before and after ageing has been reported as 0.8 to 1.2",, [5.5*]. It is also known that titanium additions lower the fracture toughness of maraging steels (Fig. 5.3). Molybdenum lowers the diffusion coefficients of a number of elements in the grain boundaries and thus reduces preferential precipitation of second-phase particles during ageing, thereby raising the ductility and plasticity of the aged steels. However, molybdenum also tends to segregate during solidification, which again induces anisotropy of the plasticity and ductility. Aluminium leads to limited hardening of martensile: at concentrations over 0.2 to 0.3",, it lowers ductility both before and after ageing. Manganese additions lead to the formation of a martensitie structure at relatively low nickel contents, but have a detrimental effect on ductility after ageing. Silicon at concentrations above 0.1 ",;, reduces the plasticity of the steels. Some elements, while not inducing ageing of Fe-Ni martensite, increase the amount of hardening 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 solubility of molybdenum and tungsten, of titanium, and of molybdenum and titanium. The decrease in 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 precipitatine phase, which, in mm. leads 10 an increase in she number of nuclei capable of growth at a given temperature and. consequently, lo a reduction in the average distance between these particles: each of these factors affects hardening. In the same w;;y as nickel, cobalt lowers both the resistance of the lattice lo dislocation movements 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 of residual ULislcnitc following hardening. As regards chromium, this element increases the corrosion 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] lo predict the 0.2",,-otlsel yield strength in terms of ihe composition of maraging steels :

0.2",, Y.S. (in 10-' psil 15.1 (",,Co) - 28..1 <",,.V1o) - 80.1 (",,Ti| 0 2",, Y.S. (in MN m=) 104 552(",,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 the strength of lXn,|Ni maraging steels results mainly from the effects of molybdenum and cobalt and. to a lesser extent, of titanium: molybdenum contributes about 50",, and cobalt 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. As regards the lS",,Ni type, three grades were initially developed: these were arbitrarily identified 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 and molybdenum, although increasing amounts of titanium are used in these three compositions as a supplemental hardener. The excellent properties of these wrought grades prompted ihe development of a cast version of this type of steel. As a result of a statistical study of compositional variations aimed at obtaining a good compromise between tensile

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

Year Alloy Ni Co Al Ti Other Mo announced

18NK200) 17-19 8-9 3-3.5 0.05-0.15 0.15-0.25 ! l9hO 18NU25O) 17-19 7-8.5 4.6-5.2 0.05-0.15 0.3 -0.5 __ 1960 18NU3OO) 18-19 8.5-9.5 4.6-5.2 0.05-0.15 0.5 -0.8 _ 1960 lSNi(casll 17 10 4.6 0.1 0.3 — 1963 15Ni-9Co 15 9 5.0 0.7 0.7 — 1963 l2Ni-2Mn 12 4.0 0.1 0.2 2Mn I96h l8Ni(35O) 17.5-18.5 12-12.5 3.8-4.6 0.10-0.15 1.4 - 1.7 — 1968 13NK400) 13 15-lh 10 0.2 1968 8Ni(500) 8 18 14 — 0.2 — 196K . l5Ni-15Co 15 15 1.0 — 0.4 — 1971 IN-763 ** 18 15 3.0 005 0.05 0.5V 1971

In wl.'\i, bal. Fc. The steels also contain - 0.03C. -• O.ISi, - O.IMn. 0.0IS. • 0.01P. Stilt in ihe product developmental stage.

53 ti|i \l I c HM \IM\i. HH. 11 MKIM.III Mills

strength and toughness, the lN\i(castl grade was developed [?.I4]. It is interesting to note 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 interest for magnetic applications at temperatures up 10 800 I (425 C). As regards the steel Jesiaiuued l2Ni-2Mn. it IN characterized b\ a 2",,Mu addition which permits the obtention of a satisfactory nuirtensiiic >trucmre with a relatively low nickel content [5./"). Research aimed at developing maragmg steels uith higher strength led later u> a rebalanced ls",,Ni composition having a nominal 350.(100 psi (2400 MN m-1 yield strength and acceptable toughness [.\/iV]. This was achieved main!) by raising the Ti content. The increased cobalt content also contributes to strength and is effective in maintaining the 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 as to retain sufficient notched tensile strength. Reversion tendencies are reduced as a result of 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 determine whether higher strength levels could be produced by heat treatment alone in Fe-Ni maraging systems with modified compositions. As shown in Table 5.1. nickel, cobalt and molybdenum are still the principal alloying elements in these steels: however, the nickel content has been reduced whereas the cobalt and molybdenum contents have been considerably increased. Titanium is still used in miner amounts but aluminium is not added. The 13Ni(400) steel is of particular interest in that it offers a strength advantage over existing maraging steels while maintaining a reasonable degree of toughness (<;/. Fig. 5.3). Finally. a composition aimed at reducing niicrosegregation effects and designated LN-763 is under development; it is characterized by a higher cobalt content, relatively low molybdenum and 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 features of 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, the austenite 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 transformation in Fe - 9.5 to 33.2 at. "„ Ni has been developed [5.22]. The martensitic transformation temperatures are shown as a function of nickel content in the metastable equilibrium diagram of Figure 5.1. As stated earlier, the transformation exhibits a thermal hysteresis which increases with increasing nicke! content. Also, the Ms temperature decreases with increasing 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 into the equilibrium austenite and ferrite compositions, i.e., the martensite reverts to the equilibrium structures. The rate of this reversion reaction depends upon the temperature and, 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 berore the 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], 301 • qiiunchcil from 21110' |- 11 KlO't'j * quenched from 21401- (12001:') • quenched from 22SO lr (1250 C'l. FERRITE S20

M+F / CD x\ / MARTENSITE \\ \ oL 0 5 iQ 15 20 25 NICKEL CONTENT, wt.%

[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 heating which 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 has revealed that a surprising variety of transformations can take place in these alloys [5.23 to 5.25]. In fact, three distinct structures, and hence three separate morphologies, have been shown to exist for the b.c.c. y. phiise obtained on cooling the f.c.c. y phase. In 0 to 5",,Ni alloys, independently of cooling rale, and in 5 to 10"(JNi alloys for slow cooling rates, a structure comprised of equiaxed x-ferrite grains is obtained; these grains are less regular in shape than is the case for annealed metal. A; sufficiently high cooling rales martensite will form in the 5 to !0%Ni alloys [5.26]. Increasing nickel contents lower the cooling rate necessary to form martensite, and at about 10%Ni a completely martensitic structure 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 composition range from about 10 to 25%Ni. Finally, alloys containing more than 25%Ni transform below room temperature to a twinned martensitic structure. All these limits are not exact 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-Mo system on quenching. In the case of maraging-type compositions, examination of a series of Fe-7Co-5Mo-0.4Ti alloys containing various amounts of nickel revealed that lath martensite apparently forms for nickel contents up to 23%, whereas twinned martensite is formed at higher nickel contents [5.29]. The lack of visible surface shears on prepolished samples has raised doubts as to whether the 18Ni(250) steel always transforms to lath martensite [5..W], However, all the transmission microscopy evidence indicates that, with the exception of the 25%Ni steel, all maraging steels normally have lath martensite matrix 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 content of 6%, without decreasing the Ms point below room temperature [5.17]. However, because of its embrittling effect, manganese should be kept to much lower levels to retain adequate toughness.

55 I lUIALI lOMMMNi. Hll.IISlRlNl.nl STEELS

5.2.2. Factors Controlling Unit Mar I ensile Formation Ahhouch further studies are necessary to ascertain the precise conditions that determine whether lath or twinned martensitc is formed, the two factors which seem to he important 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 formation of twinned martenshe. as was pointed out in Chapter 2 (Section 2.3.5). The SFE can only be varied by adding an appropriate alloying element, whereas the Mf temperature am 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 the effect of any individual addition is not constant and depends upon the total composition of the alien. The effects of the solute elements commonly found in maraging steels on the Us and As temperatures of typical base compositions are shown in Figure 5.5. It can be seen that nickel generally lowers A/.,: in the case of Fe-Co-Mo compositions, this effect is greatest for nickel contents above 17.5",, [5.18]. Molybdenum strongly depresses the Ms temperature, especially in steels with high nickel and cobalt contents. Chromium also 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/s temperature, this element exerts opposing effects as regards lath martensite formation. Increasing amounts of titanium, niobium, vanadium and silicon in a Fe-22.5%Ni composition cause the A/,< temperature first to increase and then to decrease. The probable lowering effect of titanium on the Mx temperature of a Fe-Ni-Co-Mo maraging composition is also presented in Figure 5.5; it is likely that some interaction occurs between molybdenum 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 to a slight initial increase in the Ms temperature with little subsequent effect [5.31].

700,

18 22 " 26 1 2 3 i 5 6 0 0.5 1 15 t i 1 " 2 3 Ni CONTENT, wt.% Mo 03NTENT,wt.% Ti.NbCONTENT,wt.% V,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, at moderate 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 proved very helpful in allowing higher alloy contents to be included in the steels while still ensuring that lath martensite transformation remains possible. As regards the effect of 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 lowering effect on A/., probably also holds for maraging steels. As discussed in Chapter 2 (Section 2.3.6). the general dependence of lath martensite formation 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 better toughness than does twinned martensite. Moreover, control of the M« temperature of maraging 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% and adding 3%Cr [5.35]. This depresses the Ms temperature of the steel below room temperature, while maintaining its M,i temperature above room temperature. This type of steel therefore possesses a metastable austenitic structure which transforms to martensite on 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 following two mechanisms : — the fine, uniform precipitation of various intermetallic compounds or of austenite; as in the decomposition of other solid solutions, precipitation of stable compounds in Fe-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 hardening and softening factors. Foremost among the former is the distortion of the martensite crystal lattice resulting mainly from an increase in the volume fraction of the precipitates. The softening factors include coalescence of precipitates, possible changes in the martensite substructure, a.id isothermal formation of austenite, which has been shown to occur at temperatures below the reverse martensitic transformation point. Many investigators have noted that some dislocation rearrangement, and usually a decrease in dislocation density, take place during the early stages of ageing; these changes are presumably part of the recovery reaction within the martensitic matrix. The onset of precipitation probably prevents further dislocation motion.

5.3.1. Precipitation Reactions In recent years, a considerable amount of effort has been directed towards both identifying the phases precipitated during ageing of martensite, and determining the shape, size and distribution of these precipitates. For this purpose, three-component and more complex steels have been investigated, mainly by electron microscopy and X-ray diffraction analysis, although several additional identification techniques have been used. For example, in some cases the chemical composition of the extracted phases has been determined by microprobe analysis. Also notable in recent studies are the use of Mossbauer speetroscopy and neutron diffraction analysis, as well as the analysis of the angles between diffraction spots. The influence of ageing temperature or alloying on various physical characteristics such as lattice constant, electrical resistivity, dilatation and Young's modulus has also been studied in some detail.

57 C OKU 1 (ONIUNINU HK.HSIKE.Nf.iH STUKI.S

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

Lattice parameter-, IA) Phase Structure ivpe " i c <• a t NiiMo Orthorhombic Cu*Ti-type « 5.0M. b 4.224, < 4.448 r-NuTi Dt>;4 ordered hexagonal 2.5505 1 !O<)f>7 3.2St)9 !V;Mi- Hexagonal C!4-i>pc 4.74 ! 7.73 O.ttlJ Ke-Ti Hexagonal CI4-iypc 4.S13 7.855 1.633 T-FeMo Tetragonal sf.2118 ! 4.SI3 0.522 -r-reFi Cubic CsO-ixpe Hexagonal 4.146 j 25.78 5.432

The results of the phase-identification studies have been ably summarized in References 5.S and 5.9. A study of the tables presented in these papers reveals a number of discrepancies, even as regards the precipitaies identified in alloys of almost identical compositions. This can no doubt be explained in part by the uncertainties inherent in the techniques 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 their erystallographic structures are presented in Tabie 5.2.; The shape of the precipitates has been 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 as precipitating 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 longer axes parallel to the 111 directions of the matrix [5.41]. The orientation relationship between the NijMo orthorhombic precipitate and the b.c.c. martensite matrix has been tentatively described [5.41] as :

(0»0)Ni,Mo // (0")* ll00]NijMo /

5.357= compression

Lb%//COHla 3aa% compression

DOO]p//[iTi]a 186% expansion

' 50.000 a) atomic mislil along Fig. 5.6. — Electron micrograph of I8Ni crystal axes of precipitate !p). b) distortion in (Oil) plane of matrix (1). 3 (250) steel aged 8 hours at 900°F (480 C). Fig. 5.7.— Distorted matrix lattice surrounding N13M0 precipitate in 18Ni(3OO) raaraging steel. After J.M. CHILTON and C.J. BARTON [5.36]. After K. SHIMIZU and H. OKAMOTO 15.41].

58 5. Ni-Co-Mo 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 the matrix. Figure 5.6 shows the distribution and morphology of the precipitate in the lRNi(25O) steel aged for 8 hours at 900'F (480'C), while Figure 5.7 illustrates schematically the distorted matrix surrounding an Ni3Mo particle. It has also been suggested that the Ni.-|Mo precipitates produced after a conventional maraging heat treatment of several hours at 900°F are metastable, and that after ageing for longer times and/or at. higher temperatures these precipitates are replaced either by Fe2Mo [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, but its growth would be limited by increasing coherency stresses which would favour the nucieation of the equilibrium precipitate [5.50]. Another tendency is to identify the molybdenum-containing precipitate as an Fe^Mo Laves, [i-Fe7Moe or c-FeMo phase. Precipitation of such compounds, or at least of those of the AiB or A7Be types, is apparently favoured by lower nickel or nickel -+- cobalt contents, or higher molybdenum contents [5.IG]; this effect might be related to a lower electron : atom ratio. The c-FeMo phase has been identified in the 13Ni(400) grade aged af 900°F for 4 hours (Fig. 5.8). As regards titanium-containing precipitates, they have been identified essentially in Ti-rich compositions; Y]-NijTi is the compound quoted most often, although a cr-phase has occasionally been reported. It is also possible that some of the titanium combines with molybdenum to form a precipitate such as Ni3(Mo/Ti). The exact role of titanium during maraging of the Mo-containing ailoys is therefore rather difficult to assess. Chemical analysis of precipitates extracted from I8%Ni alloys has shown that the true compositions of the precipitates are not as simple as those represented by the formulae in 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 much moVodenum for it to be NijMo, and it has been tentatively identified by some authors as 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 occurred during maraging. Some evidence has been found that suggests that pre-precipitate zones may form during the initial stages of age hardening. From observations of diffraction streaks in a sample of the titanium-rich 25%Ni steel after cold working, refrigeration and ageing for 2 minutes at 890°F (475°C), it was suggested that there is a G.P.-zone stage in which segregates form parallel to the a directions [5.43]. Diffraction pattern streaking was also observed in the I8Ni(250) steel and some higher-titanium alloys [5.52], and the following hardening mechanism involving titanium was proposed : lath martensite —>• metastable b.c.c. ordered Ni3Ti zones (D03 structure) -v Widmanstatten precipitation.

59,, nMMMM. IIK.II SIP.lNdl H sn-HS

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-htMo compound, work on 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 dislocations or 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 normally loitnd. No d»mht the kith martcn.siie dislocalion structure is helpful in providing a mia-i-uniform distribution of sites for precipitation. As aheady indicated, the precipitates arc freo,uentl\ frehe\ed to be coherent with the matrix, and in many cases strain fields have been 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 orientation of the dislocations within the martensite matrix [5.54]. If precipitation occurs along the leneth of the dislocations, the precipitates would tend to lie in the 111 i directions already referred to [5.43].

5.3.2. Ordering As already pointed out. an ordering reaction in the cobalt-containing solid solution can contribute to the hardening of maraging steels on ageing [5.6. 5.4J. 5.32, 5.55. 5.56]. As regards the binary systems involved in this type of steel, ordering has been observed over a wide range of solid-solution compositions in h.c.c. Fe-Co alloys; it is known to occur in fee. Fe-Ni solutions at 50 and 75 at.",, Ni. whereas no ordering has been observed in fee. Co-Ni sohd solutions. On the basis of neutron diffraction experiments, it was stated that B2-lype long-range ordering developed on ageing an Fe-22.7Ni-19.3Co alloy [5.55]. However, neutron diffraction experiments have also shown the absence of long-range order in ISNi - 8 W 12Co maraging steels [5.49. 5.57]. From similar work performed on the l8Ni(35O) steJ aged 3 hours at 950 F (510 C) [5.57]. it was suggested that a high degree of short-range order is established in localized regions. Since the magnitude of the interatomic attractions decreases in the order Fe-Co. Fe-Ni and Co-Ni, Co-Fe-rich short-range ordered regions would form preferentially, the remaining regions being nickel-rich. These short-range solid-solution atomic arrangements may control the processes leading to precipitation of the Mo-Ni and Ti-Ni intermetallics in the nickel-rich regions. The diffusion of iron towards these regions being inhibited by the Co-Fe attractions, longer limes or higher temperatures are required for the replacement of the nickel-rich phases initially formed by iron-rich precipitates [5.58]. This reasoning is compatible with the metastability 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 and cobalt contents very high strengths can be achieved. Figure 5.9 gives the curves of yield strength versus cobalt content for two different Ni levels and two different maraging heat treatments [5.21]. The strength appears to increase approximately linearly with increasing Co concentration. However, one of the most important contributions to the strengthening 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 added simultaneously 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 :ncrements produced 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 the exception of the molybdenum-containing alloy, the addition of 8"/,,Co increases the yield strength by about 20.000 to 45,000 psi (140 to 310 MN/m2). This is of the order of the strength increase produced by cobalt alone, as illustrated in Figure 5.9. Thus, in these alloys, cobalt supplies only a relatively small, roughly additive strengthening contribution. In the molybdenum-containing quaternary alloys, however, the strength was increased by 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 to explain 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 addition of 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 in the martensite matrix and thus favour precipitation of molybdenum compounds [5.46, 5.52, 5.59]. Electrical-resistivity measurements have also shown that the principal effect of cobalt is to increase the amount of molybdenum precipitated. This technique proved very useful since it was found that resistivity is directly proportional to the amount of molybdenum in solid solution, while changes in cobalt content have a negligible effect on resistivity [5.60]. The measurements also revealed that the resistivity changes due to the recovery of the matrix during ageing were very minor compared with those produced by precipitation. Figure 5.10 illustrates the resistivity changes observed in Fe-18Ni-5Mo ternary and Fe-l8Ni-5Mo-8Co quaternary alloys during ageing at 850°F (455=C). Both curves show an initial loss in resistivity due to precipitation of molybdenum, followed by an upward turn due to austenite formation. The cobalt addition tends to accelerate the ageing reaction and to remove more molybdenum from solution. It has also been suggested that cobalt may alter the dislocation structure of the martensite matrix to provide more numerous and more uniformly distributed nucleation sites for

-2000

: •t0- ••• •• •'_ •-. -v-B 40 COBALT CONTENT, wt.7.- 1000 Anneale .001 .01 ; o.i l.o . AGEING TIME, hours

Fig. 5.9.— Yield strength of Fe-Ni-Mo alloys. After S. FLOREEN [5.21]. Fig. 5.10. — Electrical resistivity of two Fe-18Ni • Maraged 24 hours at 800°F (425°C) alloys as a function of ageing time at 850°F (455°C). • Maraged 3 hours at 900°F (480°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 atistenite arains occurs and the refilling inanensitic structure is therefore much finer [5.-J6]. Cobalt may also decrease the SFL of the matrix in the ausienitic state, which v\otild discourage cross-slip and retard cell growth. The mean dislocation density of marlensite would he increased, 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 cobalt would 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 during maraging. and the kinetics of these can be studied reasonably well before reversion of the martensite matrix to austenite becomes a major factor. Firstly, several authors have observed some dislocation rearrangement and loss during the initial stages of ageing. Internal-friciion experiments [5.62] have led to the conclusion that this recovery reaction is 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 an Fe-l8Ni-SCo-5Mo alloy on ageing at 850 F (455 Q. It can be seen that the hardness of both alloys rises steeply after ageing for as little as 1 to 2 minutes. As slated in the previous section, small but real changes in the resistivity of the same alloys are noted after even the shortest ageing times (Fig. 5.10). From such results, it has been concluded that the incubation lime for age-hardening is zero [5.6U. 5.62]. In many cases the isothermal ageing kinetics can be conveniently described by a relationship of the type : _J.Y v,, — Kt". where .v is either the hardness or the electrical resistivity at time /. .v,, is the value of this property in the annealed condition, and K and n are constants [5.56. 5.60. 5.63]. Quite often the values of the time constant n are of the order of 0.2-0.4. i.e.. distinctly less than the n value of 0.5 for the idealized case of diffusion-

Ternary Quaternary Coherent precipitation $ II No coherent precipitation O D

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

.01 0.1 1.0 ». 15 20 25 30 AGEING TIME, hours at.% Nior Ni+Co

Fig. 5.11. — Hardness of :wo Fe-18Ni alloys as Fig. 5.12. — Compositional limits for occurrence of cohe- a function of ageing time at 850 F (455"C). rent precipitation in ternary Fe-Ni-Mo and quaternary After D.T. PETERS and C.R. CUPP [5.60]. Fe-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, the nominal values of the activation energies are fairly lov., being typically of the order of 30-50 kcal/molc, which is well below those commonly observed for substitutjonal-element diffusion in ferrite. On the basis of these results, it was generally concluded that the maraging kinetics are primarily determined by the martensite matrix structure. The absence of an incubation period was attributed to the elimination of the free-energy barrier lo nuclcation under conditions of high supersaturation and precipitation at dislocations [5.60]. The low values of the time constant n and the activation energy are commonly interpreted in terms of pipe diffusion through the dislocations present at a high level of density in the lath manensite matrix. Evidence of the migration of substitutional solute atoms on dislocation sites in martensite emerges from (he observation of serrations on the s'ress-strain curves for Fe-Cr-Ni, Fe-Cr-Co and Fe-Co-Ni martensite alloys tested at the usual strain rates C of 10-3 t0 io-6 s-i an(j at temperatures ii; the 66O-75O F (350-400°C) range over which ageing may possibly occur [5.64]. However, other studies suggest that this interpretation is not completely correct. The ageing kinetics of an Fe-8Ni-13Mo alloy in the ferritic, cold-rolled ferritic, or martensitic condition [5.65] were found to exhibit noticeable incubation times. Furthermore, the activation eneny values for age-hardening of all three structures were of the order of 65 kcal/mole. Similarly, a value of 60 kcal/mole was determined in the case of a martensitic stainless steel [5.56]. These results show that activation energies of the order of those for diffusion of substitutional alloying elements can be obtained during ageing in a lath martensite matrix. In order to explain the rapid hardening of 18%Ni-type maraging alloys, it was accordingly suggested that the absence of incubation times is possibly related to the tendency of the higher-nickel alloys to form A3B precipitates and to the close fit with the b.c.c. matrix that is generally possible with precipitates of this type [5.8]. The temperature dependence of nucleation may also be responsible for the low activation energies observed. As regards the effect of ageing temperature, it has been observed [5.60] that, contrary to the behaviour established for the Fe-18Ni-5Mo alloy, there is a pronounced discontinuity in 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 or resistivity values, and in the peak hardness vs. ageing temperature curves. More recent studies [5.28] on Fe-Ni-Mo and Fe-Ni-Co-Mo alloys have shown that two precipitation processes are operative on ageing the quaternary alloys. The first takes place within the matrix and is predominant when the ageing temperature is below 840GF (450°C), while the second occurs preferentially on dislocations, predominating when ageing is performed above 840T. This complex precipitation behaviour had already been identified on the basis of electrical resistivity measurements [5.67]. The hardening that results from holding at temperatures below 840°F is probably due to the formation of ordered precipitates, 10 to 50 A in diameter, that are coherent with the martensitic matrix. They appear to have a hexagonal structure with a = 7.02 A and c = 2.48 A, and the following orientation relationships with the matrix :

[100]p // [211]a [011]P // [lll]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 of coherent precipitation appears to foilow approximately an electron concentration contour which corresponds to an average number of paired electrons per atom, epja, of about 6.0. The formation of fine, coherent precipitates of an ordered phase on ageing the 13Ni(400) maraging steel at temperatures between 750 and 840°F (400 and 450°C) has

63 UHIM I UIMMMM. MICH SIKKNCilH SHUIS

M HJectron microdifTryciion. show- tii.--.irun micrograph. 40.1HK) 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 hours at 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 clustering is 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 supersaturation of molybdenum in the presence of cobalt [5.60]. However, the similarity of the results 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 particular process [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 alloys containing chromium, nickel, cobalt and molybdenum suggest that the low-temperature ageing reactions in maraging steels are governed primarily by the formation of Ni- and Co-rich zones, and that the chromium and molybdenum only participate in an auxiliary manner. As regards the second mode of precipitation, i.e.. that occurring on dislocations on holding at temperatures above 840 F. it ieads to relatively faster hardening of the martensite; the presence of cobalt in the quaternary alloys accelerates the precipitation reaction in the same way as does raising the molybdenum content, and confirms that the former element increases the supersaturation of the latter in these alloys as compared with ternary alloys [5.28].

5.4. Austenite Reversion

As stated previously (

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'cct of an individual alloying element on reversion may be partly explained on the basis of whether it tends to stabilize the anstcnite or the ferrile phase [5.75]. However, a much stronger effect probably results from ihe change in matrix composition thai accompanies precipitation. Titanium, for example, has been found to retard reversion markedly, due to 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 formation of Fe2Mo [5.6.?, 5.74]. The effect of cobalt and molybdenum on reversion in maraging steels during tempering has 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 steps are required in the more complex alloys. As indicated in Figure 5.14, the initial rnartensite first 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 nickel contents (Fe-Ni-Co alloys). Therefore, it is apparent that cobalt and molybdenum additions 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 is predominant. Austenite usually starts to form at the martensite platelet boundaries; once the reaction has progressed far enough, the structure is lamellar in appearance, with elongated austenile ribbons strung out along the boundaries, as shown in Figure 5.15. In addition, smaller austenite pools form within the platelets. The precipitated particles have occasionally been observed to dissolve in the austenite. With prolonged holding at temperatures of the order of 750-1300'F (400-700X), very appreciable amounts of stable austenite can be formed. For example, maximum amounts of 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 be transformed to martensite by refrigeration. Cold working, however, will cause it to transform [5.79]. The results of a recent investigation [5.80] suggest that austenite

SYMBOLS 1st step MARTENSITE -* Ctp + Yr a = ferrite Co

Y = austenite \ on tooling. Yr—Mr M = martensite C = nickel content Z Fe-Ni-Mo Ms = martensite trans- F*-NI-Co-Mo Fe-NI-Co formation point 7 SUBSCRIPTS 2nd_stet> r. p - Nl-depleted

Fig. 5.14. — Schematic diagram showing possible paths of mar- Fig. 5.15. — Electron micrograph of tensite-to-austenite transformation in Fe - 20Ni - 3 to 15Cc, 13Ni(400) maraging steel agsd 4 hours Fe - 20Ni - 4 to 8Mo and Fe - 20Ni - 5Mo - 2 to 9Co alloys. at 1110°F (600°C) and air cooled. After C. SERVANT and G. CIZERON [5.77]. After A. MAGNEE et at. [5.7]. x 30.000 i >U \1 , i i >\ I \IM\i. Mil .11 si XI M. Ml Sill 1 S

^ lii;. 5 U".-- LlTtvloftcmpi.'1'.iiBtcrnpcrii- ^-'^ uiredinic; ~hlon saturation in.i^ncli/a- i" non and C urie point ol 1SNI<-\^\i.'rageing on stress-corrosion beh'.ivio r is not clear-cut, formation of large amounts of reverted ausienite during an overagetng treatment seems to increase the stress-corrosion resistance of maraging steels [-5.-/-J. The magnetic properties of maraging steels are also afieeted by the presence of reverted austenite [5.16. 5.50. 5/>J. 5.74. 5.79]. As an illustration. Figure 5.16 shows the effect, of tempering temperature on the saturation magnetization and Curie temperature of the lis'Ni(250) steel. As austenite starts to form, the saturation magnetization \the value of which depends inversely on the amount of austenite present) decreases appreciably. Simultaneously, there is a marked increase in coercive force. This is consistent with the fact that large changes in coercive force arc to he expected only when particles of the dimensions of the domain walls, i.e. ~ 1000 A in size, are present. Precipitate particles are generally too small to affect ;he properties significantly, but the austenite panicles formed by the reversion reaction are of the right size. Austenite can also he formed by heating an alloy above its ,-J., temperature, but in this case 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 austenite contains a high density of dislocations [5.81 to 5.83]. Orientation relationships indicate that 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 the martensite matrix to austenite • ferrite by the diffusion-controlled reaction may occur. When this happens, the subsequent shear reaction may be affected, apparently due to the fact 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 particles formed during heating, and/or to the corresponding decrease in the nickel content of the martensite. When additional alloying elements are present, the changes in the shear reaction with composition and heating rate can become very complex indeed. In the case of the I8%Ni steels the general trend still appears to be that decreasing the heating

66 \i (o-Mu \I\K\(,|\(, Sill IS PHVSK -\l. MLIALLL'kGY

rate raises the .-(., temperature [5.W]. but the multiplicity of reactions makes inlerpretation of ihe data difficult. As the temperature is increased well above .),<. a homogeni/ation reaction becomes predominant. /.<•.. precipitales dissolve and concentration gradients due to precipitation or reversion are removed [5.7-/|. The lime and temperature required for complete homog.-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 and toughness. As discussed in Chapter 2 (Section 2.4.3). approximately one third to one half of the yield strength of fully heat-treated IS",,Ni maraging steels can be ascribed to the strength of the Fe-Ni lath martensite formed on cooling from the aiineaiiim temperature, the rest being due to the line dispersion of precipitated particles which form on ageing [5.9]. It was also shown that about three-quarters of the overall strength of Fe-Ni lath martensitcs stems from solid-solution hardening, and the remainder is due to 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. It has often been noted thai elimination of carbon and other deleterious impurities should be beneficial in this respect, as should the relatively uniform precipitate distribution achieved in a lath martensite matrix by age-haraening. Molybdenum appears to play an 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 of 2",,Mo [5.5V]. The cfleets of varying the nickel, cobalt, molybdenum and titanium contents on mechanical properties and microsegregation in maraging steels were recently examined [5.2/]. A nickel content of i8",, gave the best mechanical properties. Molybdenum additions raised the strength and gave substantially better impaci values at all strength levels. Increasing the cobalt level from 12 to 20 "„ gave continuous increases in strength and decreases in toughness in Fe-l8Ni-Co base alloys. Good tensile and impact properties associated with reduced 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 from 0.4 to 0.1%. The possible role of the high nickel content of the matrix in preventing fracture must also be considered. In low-alloy steels, nickel reduces the tendency to cleavage and lowers the ductile/brittle transition temperature. In maraging steels, this effect of nickel could be helpful in minimizing hydrogen embrittlement, where quasi-cleavage fractures have generally been observed. Normally, however, fracture in maraging and other high-strength steels occurs by means of void nucleation giving rise to localized ductile fracture, and it is 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 impurity contents and a lath martensite matrix with a fairly uniform precipitate distribution do 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 fracture toughness of the 18Ni(25O) grade is significantly higher than that of the other grades compared at equal strength levels [5.59]. Since the 18Ni(250^ grade is age-hardened by a Mo-rich precipitate which is metastable and eventually £jes into solution, it has been suggested [5.5] that this process may start during maraging to give a structure of Ni3Mo particles with adjoining austenite that might be more effective in preventing void formation at the precipitates, and thus in delaying fracture.

67 I i IBM I 1(

,,. NWo-Mo MAR.\C;IM; sniis — 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 5 t Table 5.M will be included where appropriate. This will apply in particular to the cast grade (Sections (>. I. (v2.I, d.2.41. the heat-resisimg grade (Section 6.2.21 and the magnetic grade (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 in a consumable-electrode furnace decreases segregation and improves cleanliness, leading to 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 performed m air ['>./]. Vacuum-induction melting of this steel has been recommended to prepare investment castings with good surface tinish [6.2. 6..*]. Finally, a recent study on the ISNi(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 such as rolling, forging (including drop forging!, drawing and extrusion. Prior homogenization b> soaking the ingots at about 2300 F (1260 C) is recommended. Hot working may be c.irried out between 2300 and 1500 F (1260 and 815 Ci. but scaling is minimized if the siarting temperature is limited to 2100 to 1900 F (1150 to 1040 Cl. Finishing at low temperature i- desirable for the obtention of uniform smaii grains and optimum mechanical properties. These steels have been shown to be self healing and to forge-weld like carbon 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 sections and for 1 hour per inch for heavier sections, followed by air coofing [6.5]. At this stage the material consists of a soft, low-carbon, high-nickel martensite which is readily amenable to machining. Since the martensitic transformation is independent of the cooling rate, hardenability is not a problem. Maximum strengthening is obtained on ageing 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 at 1800 F (980 C, for 4 hours, overageing at 1 !00'F (595=C) for 4 hours, solution annealing at 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 maraging grades 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 in Table 6.1. It is seen that the wrought and cast grades have good combinations of strength, 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, while its room-temperature Charpy V-notch toughness is higher; this improvement is attributed to the fact that use of segregation-free spherical powders produced by the rotating electrode method has eliminated banding [6.4].

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

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

L! IS.. 0.2' ,. VS.. El. R. A.. C\ 'N K .-. N. T.S. Hard- Moduli. Pois- Material (2 in.). impact |Q« j in. ness. <7 v son's Ref. 1 IO psi MWm- 10' psi fi.lh ./ A/.Vni -! - L. T.S. R. IO"psl — ratio

ISNiCOO) 195-230 1345-15X5 190-225 1310-1550 6-12 35-67 26-50 35-6X 100-160 110 176 1.5 44-68 26.2 1X1 0.26 [6.5] bar' (A 10) (elitsticin } ISNiCSO) 245-270 169(1. ,1M) 240-265 1655-1X25 6-10 35-60 18-33 24-45 90-150 99-165 1.5 4K-5O 27.0 1X6 0.3 \6.5] bar' (A 'Ol (elasuein I 10.4 72 (rigidity) ISNi(3OO) 265-305 1X25-2105 260-300 1790-2070 5-10 30-50 12-19 16-26 80 130 US-143 1.5 51-55 27.5 190 0.3 [6.5] bar" (A'i 10) (elasticity! 1 Casi grade' 249 17IX 237 1635 10 40 i5 20 - - - [6.3]

1 P M bar' 282 1946 27S I9IX 10 54 23 31 - 5" — - - V>-4] '"

Cuniiiiion : 1500UF (SIS' Ci - I h 900cF (480^CI - 3 h. l imdilion : 1800:K ("980 C.) - 4 h ' IIOff'F (595°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 quenching: used at •Hid I 1480 Ci tor 3 h.

In Figure 6.1 the fracture toughness of the three basic maraging grades is compared wilh that of M.eels of similar strength. The superiority of '.he former is immediately obvious. It has been shown [6.9] that the puriiy of the raw materials used strongly affects the toughness properties of the steels. In particular, the sulphur, carbon and nitrogen levels should be carefully controlled in order to prevent harmful grain-boundary precipitation of intermctallic c impounds [6.10]. This type of grain-boundary embriitlement can be reduced 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 less than 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 been set at 0.005";,, while the simultaneous presence of siii-on and manganese at the 0.15°,,

ULTIMATE TENSILE STRENGTH, MN/m* 1250 1500 1750 2000

Fig. 6.1. — Effect of strength level on fracture toughness ol_ ofrepresentative high-strength 200 250 300 steels. After J.C. HAMAKER ULTIMATE TENSILE STRENGTH, 103 psi and A.M. BAYER [6.8].

69 t•ifflALT-l CM AININii HIliMSTRf-MiTH STFEl S

20 r -,2200

.300!—

280

40 50 40 " 50 COLD WORK V.COID WORK

Fiu. h.2. — Effect of cold work and maraging parameters on strength and fracture toughness of 18Ni(25O) crade. After E.P. GILEWICZ [6.18], [nilial condition (0.375 in. plate) : 1500 F (815 Cl - 1 h, AC.

330

30 40 50 60 V.COLD WORK

Fig. 6.3. — Effect of cold work and maraging parameters on strength and fracture 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 maraging steels 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 in amounts of 0.5",, appear to be harmless [6,14]. Vacuum-consumable electrode rcmeliing has been found [6.15] to increase significantly' the short transverse strength of the I8NH2OO) 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 some evidence 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 showed that maximum ageing response was obtained on treating at 90CPF (480 C) (see. for instance, Ref. 6.IS), and this temperature was thus adopted as the standard ageing temperature. However, there are cases in which departure from the standard practice may be indicated, e.g. to optimize fracture toughness. Work in this respect has mainly been concerned with the high-strength grades, since the 18Ni(200) grade exhibits excellent fracture toughness in the standard condition (see Table 6.1); very often, it was carried out on plate or sheet products and combined with an evaluation of the effect of cold work. Figures 6.2 and 6.3 shov, the effects of cold work and maraging parameters 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 to marstraining has already been dealt with in the preceding paragraph; a more complete description of the effect of cold work on the mechanical properties of the 18Ni(25O) and l8Ni(300) grades is provided by Figure 6.5.

(12% OFFSET YIELD STRENGTH, WN/m? (200) nJgOO 150D • •••: - 2000- • - •yV,-. -21DU.-":— : V 2200 2300 '••-.•:,-...-. , 771. T . • : |W 2Q.°/.C.W.«S00°F(4B0°C Hi ^v^900»F(4BO°C|-2h

••••- :/•]'•: 120H — 40% C.W..900 °F(4B0'O3h 70%C.W. 025% CW»900°Ft480°C)-9h .900°FKB0°C)-3K!20 tgioo "^i_JpiJ°F(«25 °C)-4h B0 = 7SO"F(400°C)-4h ! TO-1

S00°n425tC)-30Dh

—'60 75%CW..900°F°||

40 260 270 1b0 29D 300 310 320 330 340 02% OFFSET YIELD STRENGTH 103 psi

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

Initial condition : • 0.06 in. sheet, 1500°F (8!5°C) - 15min, A.C. ] ., tjtuj:nai) After H VV MAV- o 0.08, 0.12 or 0.24 in. plate cold rolled to LR°Jr 'andCC BUSCH' [6 '9] 0.06 in. sheet (25, 50, and 15% reductions) ) NOR' Jr" and 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 COBALT-CONTAINING HIGH-STRENGTH STEELS

2300

2200

2100 300 T LONGITUDUWL LONGITUDINAL - 2000 1 __-. TRANSVERSE

20 60 70 20 30 50 60 %CQLD REDUCTION

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) and iNNii,300) grades are shown in Figure 6.6. The tensile properties of the 15Ni-9Co grade developed for use at 10004F (540°C) have been included for comparison. The latter's stress-rupture properties are given in Figure 6.7, together with three data points for she ISNi(250) and lSNi(300) grades. The i5Ni-15Co grade recently developed for use at moderately high temperatures as a high-strength soft-magnetic material also exhibits good 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 measured from the slope of the creep curve at the 3500 hour point) at 850°F (455°C) is of the order of ftO.OOO psi (420 MN/m-). As regards cryogenic properties, the results of an investigation on I8Ni(200) plate are shown in Figure 6.8«. Although the A'u. value was considerably lower at —320°F I 19(vC) than at higher temperatures, this stea.l exhibited an excellent combination of strength and toughness at all temperatures. The variations in the ultimate tensile strength of l8Ni(?5O) and l8Ni(30O) sheet are shown n Figure 6.86. A nominal strength of 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, fully martensitic 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- cationstin .

Bf-liarhur Representative fatigue properties for the three grades, obtained from rotating beam tests oil ;uod bars* are given in Figure 6.9. The endurance limit increases with the tensile strength of Hie grede. Smooth ters of the !8Ni{200), 18Ni(250) and 18Ni(3OO) 6. Ni-Co-Mo MARAGING STEELS — THE CONVENTIONAL GRADES

Fig. 6.6. — Elevated-iemperature properties T5ST TEMPERATURE,°C of 18NK250), 18Ni(300) and 15Ni-9Co (high- 0 100 200 300 400 500 temperature) maraging grades. After [6.6] for — 2000 250 end 300 grades, and S. FLOREEI and R.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 of the 15Ni-9Co high-temperature grade. After S. FLOREEN and R.F. DECKL'K [6.2/]. Data for 18NK25O) and !8Ni(3OO) grades inclu^ej for comparison are from [6.6].

200 400 600 eoo BOO 30 W 300 TEST TEMPERATURE.°F RUPTURE LIFE, hours

TEST TEMPERATURE, °C TEST TEMPERATURE,°C -200 -150 -100 ~ -50 -200 -100 230

a e85"F(475"C)/8h -\ o 900°F(480'C)/3h

60L -30B -200 -100 0 -320 -200 -WO 0 »70 -ma TEST TEMPERATURE. °F TEST TEMPERATURE. °F a) 18Ni(200) plate. After C. VISHNEVSKY and E.A. STCI- h) lSNi(25O) and (8Ni(300) shecl. After D.L. CORN [6.24]. Condition: 150DoF '815 C) - 1 h. OERWALD 16.23]. Condition: 1650°F(900°C) -2h. A.C.; c c 1450cF (790°O-2h. A.C.; 900°F t480°C) - 2 h. A.C.; 90O°F (4S0 O - 3 h or 886^ (47« O - 8 h. Fig. 6.8. — Low-temperature properties of the 18Ni(200), 18Ni(250) and 18Ni(300) grades.

73 COBALT-CONTAINING HIGH-STRENGTH STEELS

24U —1600 ea ieNi(300) 20a SO 18NK250) ^ 18Ni(200) - 1200 .160 i 1 s n SOU in "• in 80 ij.--— tn I ton 40 Fig. 6.9. — Typical rotating-beam faligue properties of the l8Ni(200), (1 1 II i 1 III 1 1 l i 1 I 1 1 n 1 7 18Ni(25O)and 18Ni(300)grades[6.5]. 10 n» n NUMBER OF CYCLES TO FAILURE 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 to 130,000 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 on the high side of each of these ranges. As regards the fatigue crack propagation rate, it has 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 limit of the three grades [6.26]. Similarly, nitriding 18Ni(250) parts has been shown to improve their fatigue strength by some 20,000 psi (140 MN/m2) [6.27]. Finally, a comparison of the fatigue data for maraging steels and low-alloy steels of equivalent tensiie strength [6.28] has indicated that both groups are equally affected by the presence or absence of compressive surface stresses, and that their endurance limits are comparable, provided the parts are machined after ageing. Surface hardening by shot peening also greatly enhances the fatigue strength of the maraging grades [6.26, 6.28].

6.2.4. Stress-Corrosion Characteristics An exhaustive review paper on the stress-corrosion and hydrogen embrittlement behaviour of maraging steels was prepared recently [6.29]. It is generally recognized that the 18";7Ni maraging steels compare favourably with other high-strength steels and offer as good, or better, threshold plane-strain stress intensity values (KUcc) over a wide range of strengths. In terms of critical crack size, maraging steels can tolerate larger 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-corrosion resistance. Marginal improvements were obtained by raising the carbon content of an l8Ni(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 was decreased to 30 ppm or less. Of the possible desulphurizing methods, electroslag remelting appears highly promising, sulphur contents of less than 10 ppm having been obtained using high sulphur-capacity lime-fluorspar slags [6.30]. As regards the effect of heat treatment, a survey of available data on the effect of annealing temperature has shown [6.29] that response to this parameter is not entirely

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

• 800^25 °G)/10h A 90Q?rTft80°C)/100h o 900°F(A86°e)/3V2h 50 jf

UJ 20 c in 10 <

2 I I 1 I 100Q 15000 100 1000 10,000 TIME TO FAILURE, minutes

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, it should 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 homogenization, overageing and annealing steps mentioned in Section 6.1 has resulted in a marked improvement of the stress-corrosion resistance in a 3.5%NaCI solution; this improvement has been attributed to grain refinement [6.3]. On the other hand, there is 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 two deaerated 3%NaCl solutions with different pH values; it is seen that time to failure is relatively 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 two NaCl 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; in 3.5%NaCl solutions, a A'lscc value of 45,000 psi y in. (49 MNm-3/2) has been reported for 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%NaCl on three experimental grades with increasing Ti contents have confirmed the dependence of stress-corrosion susceptibility on yield strength : the steel with 0.65 %Ti (analogous to n 18Ni(300) : Y.S. 293,000 psi, i.e., 2010 MN/m2) had a KUcc value of 11,000 psi y > - (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'in. (51.7 MNm-3/2) [6j4]_

In more practical terms, it has been stated [6.35] that the 18%Ni rnaraging steels can be used at high stress levels in non-corrosive atmospheres (e.g., moist air) or at medium stress 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 yield strength. 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 COBALT-CONTAINING HIGH-STRENGTH STEELS

yield strength. Finally, their stress-corrosion susceptibility at high stress levels in NaCl solutions saturated with H2S (/>H = 5) is large. As regards the mechanism of stress-corrosion cracking in maraging steels, the literature supplies a considerable amount of indirect evidence favouring a hydrogen embrittlement mechanism rather than an active-path corrosion mechanism [6.29]. Concerning this type of embriulement, many " direct" studies have been carried out on maraging steels either precharged with hydrogen [e.g. 6.34] or tested in hydrogen gas. It has been shown that, although these steels are subject to hydrogen embriulement. they tolerate greater quantities of hydrogen than other high-strength steels for a given loss in ductility or a given susceptibility to delayed fa lure; furthermore, they recover their original ductility much 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 harden slowly [6.5, 6.6]. They can be reduced by substantial amounts before intermediate annealing is required; the reductions achieved are of 75% or more in cold rolling, of up to 85 "a in wire drawing, and of 30 to 40% in deep drawing. The standard intermediate annealing temperature is 1500°F (815°C), with times of 1 hour per inch of thickness; provided a 30% cold reduction is performed in the final operation, this temperature can be raised to 1800°F (980°C). In addition to the three cold-working operations mentioned above, the maraging steels can be fabricated by processes such as tube spinning, shear forming, explosive forming, hydroforming, bending and shearing. It should be noted that even though their work-hardening rate is low, the annealed ductility as measured by elongation in 2 in. is comparatively low, ranging from about 10 to 25"O. This is of particular importance in forming applications involving an appreciable amount of stretching; in the forming of deep-drawn shells or in cupping operations, drawing of the material rather than stretching should be the basis of tool design. However, stretchability has been found to be significantly improved by a prior overageing 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 makes them ideally suited to cold-heading operations. Two recent review papers are available on the machining [6.38] and welding [6.39] of maraging steels. These steels are machined most easily in the solution-annealed condition, in which the parts can be finished to their final dimensions since dimensional changes and distortion on ageing are minimal. Their machinability after ageing is comparable to that of conventional steels of similar hardness. Grindability is essentially the same as that of ordinary constructional steels, provided a heavy-duty water-soluble grinding fluid is employed. The maraging steels can be torch cut; plasma arc cutting is a 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 can be 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 properties in the heat-affected zone and to develop good strength in the we •! metal. Gas-shielded processes are generally favoured. TIG welding offers virtually no problems; on the other hand, MIG welding occasionally produces porosity, but this can be eliminate^ by closer control of welding parameters or filler composition. Short-arc welding requires the use of pure helium shielding to produce welds with good mechanical properties while maintaining adequate operability [6.5]. Details on filler metal compositions, processing parameters and other joining methods will be found in the review paper mentioned above. Joint efficiencies of 90 to 95% are readily obtained with good weld ductility 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 applications of the conventional maraging steels can be grouped under three headings : — aerospace, aeronautic and marine applications : these include rockel motor cases with diameters ranging from 9 to 260 inches, and the load cells capable of measuring their thrust; pivots for the gimbal support of a missile trans-stage engine: the torsion-bur suspension system for the Lunar Rover Vehicle; flexible drive shafts for helicopters; aircraft landing-gear components; hinges for swing-wing planes; foil assemblies for hydrofoil ships; deep-submergence marine vehicles; — structural and machine components : pressure vessels, components of liming mechanism in fuel injection pumps, index plates for machine tools; bolls and fasteners; barrels for rapid-firing guns; parts operating at cryogenic temperatures; springs and .estraining devices in miniature instrumentation; — tooling applications : die-casting dies or components thereof; moulds for the plastic industry; rams for extruding lead or lead-tin sheaths; cold-forming dies. This great diversity of examples reflects both the excellent combination of properties and the very high reliability of the steels. It has been shown [6.41] that the reduced manufacturing costs associated with their remarkable fabricability can more than compensaie for the difference in price between the maraging steels and low-alloy steels of equivalent strength. These factors should favour the development of new applications in 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 has progressed 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 induction melting and consumable-electrode vacuum remelting [/ /]. The same process was used successfully 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; it was 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"s casting and hot-deformation behaviour is similar to that of the conventional maraging grades, 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 temperature commonly used for the conventional maraging steels [7.1, 7.3]. The recommended heat treatment for the 18Ni(35O) grade is as follows : annealing for 1 hour at 1500T (815°C) .ollowed by ageing at 900°F (480°C) for 12 hours or preferably at 950°F (510°C) [7.1, 7.4] or 1000°F (540aC) [7.3] for 3 hours; lowering the annealing temperature from 1500 to 1450°F (815 to 790°C) has also been recommended when prior forging is carried out at !900°F (1040C) instead of 2100°F (1 i50°C) [7.3]. The steel also shows an interesting response to ausageing [7.7]. As regards the 13Ni(400) grade, a recent investigation [7.5] has shown that homogenization at 1800°F (980°C) for 1 hour followed

77 COBALT-CONTAINING HIGH-STRENGTH STEELS

by ageing for 4 hours at 980°F (525°C) appears to produce the best compromise between strength and toughness. Informavion on other aspects of processing is restricted to the 18Ni(350) grade. Cold rolling 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 fracture toughness were obtained, but the joint efficiency was rather lower than that achieved with the conventional maraging grades [7./]. Machinability in the annealed condition is excellent, and the machined parts are expected to undergo, upon full hardening, a uniform 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 powders was 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

Condition U.T.S., 0.2 % Y.S., EL, R.A.,1 CVN N.T.S. Hard- Moduli, Form (all steps followed impact, 3 ness, CN Ref. I0 psi\/'in. U.T.S. by air cooling) ft.lb / MNm-'l' Re I 18Ni(350) grade I 0.5 in. plate (long.) 1500;F(815'O- 1 h 345 2379 331 2282 8.9 37.2 7.2 9.8 —— 0.58 58 [7.4] - 950F(510°C) - 3 h. ft = 7.5) •28.1 194 (elasticity) 1 in. bar 150ODF(815eC)-l h 358 2468 352 2427 43 12 16.3 — 0.66 17.1) ; K, > 12) 10.7 74 - 900 F (48O"C) - 3 h. .(rigidity) Ditto 329 2268 316 2179 24 T-- 9.5 {7.4) 1/4 in. bar ' vac. m. air melt Ditto 320 2206 308 2124 38 10' 13.6 « 0.2 in. ho. .ig. Ditto 332 2289 326 2248 5 38 41.8 0.54 [7.1]

rolled sheet | transv. Ditto 346 2386 335 2310 4.5 1 0.49 « (.JCt = 18) as drawn" 224 1544 182 1255 0.86 [7./] 6 0.032 in. wire as drawn 367 2530 287 1979 1.2 — « drawn11 and aged 429 2958 367 2530 1.2 « (900JF, 3 h) Pr;alloyed pow- 2000^(1095=0-2 h 332 2289 320 2206 16.5 28 38 — U.6] der extruded bar -f- 950°F(510'C) - 3 h.

13Ni(400) grade Hot rolled 405 2792 395 2724 25 0.73 61 30 207 [7.2] -f 900=F (480°C) - 4 h. (ATt=3.5) (elasticity) '/2 in. bar i !8OO=F(98O°C)-1 h 390 26*9 380 2620 27 — — 59 « - 900'F (480°C) - 4 h.

Hot rolled 368.7 2542 362.3 2498 5.2 1.5 25 27.5 — 62 [7.5] + 900°F (480°C) - 4 h. Hot rolled 256.9 1771 221 1526 4.3 5.5 60.8 66.S — 51 - lli0oF(6u0°C)-4h. "/loin, plate 18O0°F(980°C)-lh 350.6 H17 350.6 2417 0.63 0.7 44.3 48.7 — 63 + 900°F (480°C) - 4 '„. 180O°F(980°C)-l h 372.6 2569 367.0 2530 1.8 6.4 55.0 60.4 — 64 + 980° F (525CC) - 4 h.

» 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 lo the 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 heal treatment; unfortunately, the associated N.T.S./U.T.S. ratio is well below unity, in contrast to the lower-strength maraging grades. The fracture toughness of this steel has formed the subject of two major investigations [7.4 and 7.7]. It has been shown that the Ku. value can be increased from about 36,000 to 40,000 psiv in. (40 to 44MNm~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 by applying a reversion treatment [7./]. For instance, the yield strength of sheet treated for 1 hour at !500°F (815°C) + 24 hours at 1050°F (565°C) increases from 141,000 to 277,000 psi (970 to 1910 MN/m2) on cold rolling to 25% reduction. Similarly, inserting a reversion treatment of 1 hour at 1150°F (620°C) between two cold-drawing operations produces wire with a tensile strength of 367,000 psi (2530 MN/m2), associated with an elongation of 1.2%; the excellent ductility is retained after ageing, while the strength is 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 laboratory heats [7.2, 7.5] was rather low. As stated above, the optimum compromise appears to be 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, the l3Ni(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, "C 200 tOO

200 <00 200 400 600 300 TEST tEMPERATURE^F TcST TEMPERATURE, 't

Fig. 7.1. — Temperature dependence of tensile properties and impact strength of 18Ni(35O) grade After 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 COBALT-CONTAINING HIGH-STRENGTH STEELS

HQLDIN'3 A\G TEST TEMPERATURE.'C f ill

LABORATORY AlR

800 'F/Bh *3 T a

3.5Y. NaCI SOLUTION BOO'F/Bh O • ® 900'F/Bh a A » 200 h HOLDING 95O'F/3h • • O 0 200 «ffl 600 600 OOO n ioo BOO HOLDING AN0 TEST TEMPERATURE,*F TE5T DURATION, minutes

Fig. ~.2. — Effect of test temperature, with or with- Fig. 7.3. — Effect of ageing treatment on sub-critical crack growth out prior holding, on tensile properties of 13NK4O0) resistance of 18Ni(350) steel in laboratory air and 3.5% NaCI solution. crade. After WJ. BotscH and T.W. COWAN [7.2). After C.S. CARTER [7.7]. Condition O!z in. ban : hot rolled - 900"F - 4 h. A.C. Initial condition (double anneal) : !70(TF mS-C) -In-'- 1WF (8I5C) - I h.

about 900;F (480 C). the drastic loss of strength experienced at !000°F (540°C) being due to overageing and reversion to aastenite. The constancy of the hardness, tensile elongation and impact strength up to 900°F are worth noting. In contrast, the impact resistance 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 resistance to crack propagation, which would be an advantageous factor in tooling and die applications at moderately high temperatures [7./]. As regards the l3Ni(400) grade, a study aimed at determining its maximum service temperature [7.2] gave the results summarized in Figure 7.2. The temperature capability in tnis case lies near 800=F (425'C). Prior holding at the test temperature for 200 hours decreases the slope of the strength vs. temperature curve; the resulting gain in strength is accompanied by only a slight decrease in ductility. A similar effect was observed in 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), the 18Ni(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 determined for three ageing treatments (Fig. 7.3). The results showed that the material aged below 900°F (480'C) was very susceptible to sub-critical crack growth; in fact, premature brittle fracture occurred in unnotched tension specimens loaded at a slow strain rate in laboratory air. Such sub-critical crack growth was attributed to a hydrogen embrittlement mechanism. Ageing at 900 and 950"F (480 and 510°C) improved the threshold stress 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 the 13Ni(400) grade, the combinations of strength/stiffness and strength/toughness/formability 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 use in 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 in various fields — machine construction, aerospace and chemical industries, naval engineering, etc. — who required steels possessing good resistance to corrosion or tarnishing, associated with strengths in excess of about 140,000 psi (1000 MN/m-), and preferably approaching those 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 stainless systems. The basic prerequisites for the cttainment of a high-strength stainless steel are that it should contain sufficient chromium and exhibit a martensitic structure. As regards the chromium content, it is well established that this must be greater than 10%, or possibly 12%, if adequate corrosion resistance is to be achieved. However, an upper limit of approximately 17%Cr is imposed, mainly by the second requirement, viz. the feasibility of 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 ensure a microstructure essentially free from 8-ferrite at the austenizing temperature, and an Ms 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 control the S-ferrite content; in addition, cobalt only slightly depresses Ms and is therefore less likely than other y stabilizers to cause an excessive increase in the amount of retained austenite. In addition to the strengthening provided by the martensitic transformation, high-strength stainless steels are usually hardened through precipitation of carbides and/or intermetallic 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 exemplified by 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 MARTENSITIC AND CONTROLLED-TRANS.FORMATION STAINLESS STEELS (in wt.%, bal. Fe)

Designation C Cr Ni Mo Cu Tj Al Others Owner of trade name

Marteiisitic stainless steels AISI420 0.15 min. 13 AISI431 0.2 max. 16 2 — — — — — — AISI 616 (422) 0.23 12 0.8 1 1W, 0.25V — AISI619 0.3 11.4 0.3 2.8 — — — 0.25V — 17-4PH 0.04 16 4.3 3.3 0.25Nb Armco Steel AM-363 0.05 max. 11 4 . 1 _ Allegheny Ludlum IN-736 0.01 to 10 2 0.2 0.3 International Nickel PH 13-8 Mo 0.04 12.S 8.1 2.2 1.1 Armco Steel Custom 450 0.05 max. 15.0 6.5 0.8 1.5 — 0.7Nb CarTech Custom 455 0.03 max. 11.5 8.5 — 2.2 1.2 — O.25Nb CarTech Controlled-transfonnation steels PH 15-7 Mo 0.07 15.1 7.1 2.2 1.1 Armco Steel - 17-7 PH 0.07 17 7.1 1.1 Armco Steel AM-350 o.o; 16.5 4.3 2.8 0.1N Allegheny Ludlum AM-355 0.13 15.5 4.3 2.8 — 0.1N Allegheny Ludlum PH 14-8 Mo 0.04 14.4 8.2 2.2 i.l Armco 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 of the strength is derived from the precipitation of chromium carbides, and yield strengths as 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 strong carbide-forming elements such as tungsten, molybdenum and vanadium) have been developed to meet the current need for stainless steels with good properties at temperatures up to 1200 F (650 C) [8.1]. Although these steels can exhibit strengths of the order of 200.000 psi (1400 MN m-k they are most often treated to relatively low room-temperature strength 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 of this class are AISI 616 and 619. Several 12%Cr steels which derive their strength from the precipitation 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 another class of high-strength stainless steels. These steels, the nominal compositions of which are also listed in Table 8.1, have an austenitic microsti ncture when in the anneated condition, 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 2.3.7 respectively), no cobalt-containing versions of such steels have been commercialized. On the other hand, quite a number of cobalt-containing martensitic stainless steels hardened by means of intermetallic compounds, and possibly of alloy carbides, hav? been developed over 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 contain nickel (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%C and are thus receptive to both carbide and intermetallic-compound strengthening. All the other steels have very low carbon contents and thus depend essentially on the precipitation of intermetallic compounds for their strengthening; the predominant tendency towards

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

Year Designation * C Cr Ni Co Mo Cu Ti Nb Others Owner of trade name announced

AFC-77 0.15 14.5 13.5 5.0 0.5V Crucible 1963 AFC-260 (C50) 0.08 !5.5 2.0 13 4.3 — — 0.14 — Crucible 1967 " Alloy B " 0.16 14 1.0 13.5 5.0 0.22 1971 AM-367 0.03 max. 14 3.5 15.5 2.0 0.5 Allegheny Ludlum 1963 D. 70 0.03 max. 12 4.3 14.5 4.0 — present Al, B, Zr English Steel 1964 Pyromet X-12 0.12 10.5 — 6 4.8 i 3 0.08N . CarTech 1961 Pyromet X-15 0.01 max. i; — 20 2.9 CarTech 1967 Pyromet X-23 0.03 max. 10 7.0 10 5.5 — — — — CarTech 1973 Ultrafori 401 0.02 max. 12 8.2 5.3 7.0 0.8 B, Zr DEW 1969 Ultrafort 402 0.02 max. 12.5 7.6 5.4 4.2 0.5 0.05A1 DEW 1969 Ullrafort 403 0.02 max. 11 9 4.5 — 0.4 — 0.15A1 DEW 1971

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 and Fe-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 possible by favouring the formation of a low-carbon lath martensite rather than a high-carbon twinned one. Although, strictly speaking, the terminology " stainless maraging steel " should be restricted to the low-carbon grades, it is generally considered as applying to this family of high-strength steels as a whole, and will be used in this general sense throughout this and the following chapter. Finally, mention should be made of the existence of a number of heat-resisting steels with cobalt 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 on room-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 fully austenitic (y), a mixed austenitic-ferritic (y + a) or a fully ferritic (a) structure, depending on 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 will depend essentially on the cooling rate and the alloy element contents. Since the structural requirements for high-strength stainless steels impose that a virtually ferrite-free structure be attained on austenizing, it is appropriate to review here the effect of various alloying elements 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 as 1928 [8.2], but was only evaluated quantitatively at a much later date [S3]. Figure 8.1 shows the effect of adding up to .10%Co on the equilibrium structure of Fe-i3Cr and Fe-17Cr alloys containing 0.05 and 0.2%C. It is evident that addition of relatively small amounts of cobalt to 13%Cr steels broadens the y field without greatly depressing the a ;=i v or a ^ a+y equilibrium transformation temperatures. The latter point is of particular importance in steels intended for high-temperature applications in which transformation to austenite during service is undesirable. Figure 8.1 also shows that low-carbon 13%Cr steels are likely to form ferrite at the highest austenizing temperatures. This constituent, which is normally designated S-ferrite, is desirable in restricted amounts in controlled-transformation stainless steels, but must be avoided in straight martensitic ones because of its deleterious effect on strength. For this reason, a number of studies have been performed in order to determine quantitatively the influence of various elements on the S-ferrite content of stainless steels. Figure 8.2

83 I uH.UT-l O\l\lNt\C, UK.(I

CARBON COMTEMT. % 0 3 Oi 05 06 07 08

10' O ALLOYING ELEMENT CONTENT, % LOG TIME (seconds)

1 in - ;. 1 lieu of \,in.'U> .ill.'wni: clc- 1 iiz K..I, isiMhermal trans!t>rma!ion of ^ ferrite in mcni> i>n the •S-ferrue comcnt of a O. IC- lou-carbon Al-C-77 and in AKC-2M) after austeni/ation for t"Cr steel. After K.J. lRMxt t-( ,i/. [S.4]. 1 hour a! 2:00 F ti;05C). 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-l7Cr steel. According to this figure, cobalt is less effecti\e in suppressing ^-ferrite than either carbon or nickel, but is more so than copper or manganese. Similar data have been established for 0.lC-!"Cr-4Ni steels [X.-f]: the results can be summarized as follows : Elcmcm U".,. eM-eru C and N. 0.1",,i N C Ni to Cu Mn W Si Mo Cr V Al Change in ^-ferriie content lin absolute "„) -20 -IS —iO -h —3 -I -8 -8 II -15 • 19 -38 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 content of O.lC-lKCr steels [SJ]. In fair agreement with the above data, the following relationship was established for the equivalent nickel and chromium contents of steels containing up 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 higher values 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, it has 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 compositions such as AFC-77, because of the increasing amounts of hardening elements that have to go into solution, special treatments have been developed to decrease the ^-ferrite content without lowering the austenizing temperature. One such treatment consists in tempering for 2 hours at 1400 F (760 C) prior to austenizing at 1900 F (1040 C); this reduces the S-ferrite content from 5 to 2",, [8.7]; the reason for the effectiveness of the preliminary tempering step is related to the removal of alloying elements from the high-temperature ferrite phase through precipitation of carbides and intermetallic compounds. Another possibility 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 investigation [HULL. F.C., WelJina J.. Res. .Simp/.. 52, I93s 1 on cast complex 12-24",, Cr steels, in which cobalt was found to be 0.5% as effective as nickel in suppressing S-ferrite formation, and 0.25 times as effective as nickel in depressing M*.

84 K. STAINUSS MAR AGING SU IIS PHYSICAL METALLURGY

to 1205 C) and then to soak it for 1 hour at about 1900 F (i040 Cl before cooling to room temperature, 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 austenite grain boundaries, resuming in a line austenite grain size: after isothermal removal of the ferrite. a marked increase in austenite grain size is observed [H.H!]. When performed correctly, this two-stage treatment not only eliminates ^-ferrite. but also results in a better combination of strength and toughness through refinement of the austenite grain (see Section 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 austenized at 2200 F (1205 C) for 1 hour and then isothermally treated for increasing times at temperatures between 2200 and 1600 F (1205 and 870 C). The transformation curves for AFC-77 with the usual carbon content would be shifted to the left of those shown in the figure.

8.3. Transformation Temperatures and Structures Since high-strength stainless steels are designed so as to achieve a fully niartensitic structure on cooling to room temperature, the effect of composition on the A/.< temperature is of primary 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 of the Mf point, can markedly affect the properties obtained. Finally, the transformation of martensite to austenite on tempering or ageing (or during service in the case of steels used at high temperatures) is also of great significance. The present section is devoted to a discussion of these various aspects.

8.3.1. Manensilk Transformation The low carbon contents of the stainless maraging steels under review result in the formation of lath martensite, the occurrence of twinned martensite being rather unusual. The significance of achieving a lath martensitic structure, particularly with respect to precipitation kinetics and precipitate dispersion, on the one hand, and to toughness, on the other, was emphasized in Chapter 2 (Sections 2.3.3 and 2.3.6), and need not be reconsidered here. As a matter of fact, all the steels listed in Table 8.2 exhibit a lath martensitic structure (of the type illustrated in Chapter 2. Fig. 2.10) on quenching following austenization. However, in the higher-carbon steels such as AFC-77, limited amounts of twinned martensite are often co-present (Fig. 8.4).

Fig. 8.4. — Thin-foil iransmission electron micrograph of 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\INi • IIHiH N| R| NV , I II

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 IUMMONU [,S'..5].

I lemeiil K.mj:e Ha.se oiniposiiion I'r (1- \2 »l.",, (IO.1C.-4Nl Si 4-S l>.03L"-i:i r I'M 0- I* 0.0.iL~-IXr-4Ni Mi- l>-f (IDH'-I'l r-4Ni-15t i 1 I 0-1.1 OO.K"-i:t'r-4Ni-l5L'i

"C 2 4 S" 9 10 1? 'i <6 ALLOYING 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, the V/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 sec i 50 C sec) [S.I2]. In the subsequent discussion, the M, temperature should be understood as corresponding to cooling rates above 90 F sec.

The effect of chromium, nickel, cobalt, molybdenum and titanium on the Mn temperature of l2",,Cr steels was determined for a series of 0.03 ",,C compositions solution treated to obtain a completely austenitic structure. The results are given in Figure 8.5 together with the ranges of contents investigated. Using these data, the following relationship between M, and composition was established [8.5] : o M,( F) = 1530 -52( ,,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, the latter 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 (the temperature at which 50% austenite has transformed) and composition, for steels containing 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 0.1 to 0.7 %Ti [8.6] :

M50CF) =2842-85(%Cr)- 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 result of the interaction of this element with carbon and nitrogen.

86 S SI MM l.ss MM.' \(,l\(, Sill I s \l MI I Ml UR

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 on its '.-fleet on M,, as compared with that of nickel for which a factor of I is arbitrarily assigned. 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 In this form the> can be compared with tne original A.R.I, equation used lo develop i'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, considerine the approximations inherent in this kind of treatment. In conclusion, it clearly appears that the ell'ecl of an individual addition element on the Mt temperature can vary quite extensively with the overall composilion of the steels. This is illustrated in Table 8.3 for six major alloying elements (graphical representations of the influence of molybdenum and titanium will be found in Chapter 5. Fig. 5.5). It is seen that, 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 to be highly dependent on interaction phenomena. Finally, copper seems to raise the A/,,, temperature, at least in !4Cr-20Co steels.

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

Alloying Steel types EITect on A/, per \u.°,, element or compositions F C ReJerence

Ni Standard stainless — 110 —61 [S.I-4] 0.IC-18O —72 —1() 18.3] 12Cr —70 —39 [S.S] ~I0Cr —I5f>* —87* l«.6J 5-10Cr, 5-8.5Ni —72 —/0 in-15] 5-IOCr, 8.5-11.5Ni —36 —20 IS-15]

C. Standard stainless —76 —t: IS. 14] 4Ni —52 — 29 [S.S] Co-Ni —85* —47* [S.6] 6"Ni —41 —23 [S.I 5]

Co 0.1C-18Cr — 17 —9.5 [SJ] 12Cr-4Ni —9 —5 [H-5] lOCr-Ni —SO* —2S* IS.6]

Mo 12Cr-4Ni-15Co —65 —.« 18.5] lOCr-Co-Ni —65* — 36* [S.6] l4Cr-20Co-2Cu —S1 —J5 [S.12\

Ti 12Cr-4Ni-15Co 0 to —99 0 lo —5.5 [8.5] (OCr-Co-Ni +205* -114* [8-6]

Cu 14Cr-20Co + 13 +7 [S.I 2]

Effect on M;Q temperature.

87 I OH U I -l O\I \1M\C, HK.IISIKIMiUI SI M'l S

J1000 1 10 100 1000 20" Z1 22 "23 2t A.R.I (Ni-0.aCr»a.3Cr.0.6Mo) HOLDING TIME AT INDICATED TEMPERATURES, minutes

7 Fie. S.h. IsD'^frrnal austenite reversion in C'r-C'o and Cr-Ni steel*. Fig. S. . — tiffed of A.R.I, on austenile content and After K. BISCMIDT et til. \H.17]. ultimate tensile strength of Fc-Cr-Co-Ni-Mo steels. Alter R.L. CATON and G.N. MANIAR [8.I3\. Condition: Initial condition : 142(1 V 11 .i5() Cl - I h, A.C. 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 temperature may 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 adverse effects. However, if appreciable quantities of alloying elements such as nickel are added, 3 the As temperature may be depressed to well below 1200 F (650°C), leading to considerable difficulties during heat treatment.

The effect of various elements on the As temperature of the 0.IC-12%Cr steel referred to above is as follows [S.I6] : Element (1%) Ni Mn Co Si Mo Al V Change in A,, Gh 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 chromium steels having compositions balanced essentially with cobalt can be expected to exhibit considerable stability with respect to reversion. This is exemplified in Figure 8.6, in which the behaviours of Fe-19Cr-15Co and Fe-14Cr-8Ni martensites on holding at various temperatures 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-Co steel; conversely, higher temperatures are required to form a given amount of austenite within a specified time i,. the Fe-Cr-Co martensite. In both cases the time dependence of the reaction is evident. The stability of martensite in cobalt-containing stainless steels is also demonstrated by the fact that no austenite reversion occurs in Ultrafort 401 after exposure for 1000 hours at 800°F (425°C), and little reversion is observed after holding for 100 hours at 930°F (500°C) [8.6].

88 SI.MMISS MAKAdlNti SILI.l.S Pin Sl( At Mf TAIL .L'RC ,Y

The stability of the reverted ausienile on subsequent cooling is strongly dependent on co-..position. For instance, it has been shown [8.17] that the ausienite formed in the above-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 subsequent cooling; on the other hand, the reverted ausienite formed in the Fe-l4Cr-KNi sieel appears to be quite stable. It should be staled, however, that this difference in behaviour is partly due to the increase in the Mf temperature of the former as a result of the more extensive precipitation 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 be tolerated before any appreciable decrease in yield strength is observed [X.!8]. Amounts of similar magnitude are often accepted when working out the compositions of stainless maraging 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 high fracture toughness in martensitic steels [8.B. 8.19]. In the Fe-Cr-Co-Ni-Mo system, the amount of austenite that ensures optimum fracture toughness without impairing yield strength is comprised between 5 and 15",,; this amount normally includes both retained and reverted austenite, as produced respectively on cooling after solutioning and on ageing [8.13]. The principal parameters controlling the retained austenite content are composition and heat treatment. A quantitative assessment of the effect ot various addition elements on the retention of austenile is provided by the austenite retention index discussed in Section 8.3.1. The relationship between the A.R.I, calculated on the basis of Cqn. 8.7 and the retained austenite content in Fe-Cr-Co-Ni-Mo steels is shown in Figure 8.7. The rather large scatter in the data points has been attributed to the fact that a linear equation was used to describe a non-linear phenomenon [8.13]. This figure also shows that the optimum 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-77 is shown in Figure 8.8. It is seen that the amount of retained austenite increases with increasing austenizing temperature. The reason for this is that high austenizing temperatures increase the alloy content of the matrix by dissolving alloy carbides and intermetallic particles, thereby depressing Ms and My.

AUSTENIZING TEMPERATURE,'C_ 950 BOO 1050 1150 30

1700 1800 1900 2000 2100^ AUSTENIZING TEMPERATURE, 'F

Fig. 8.8. — Effect of austenizing temperature on retained austenite content in AFC-77 after lempering at 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 \1MNI i Uli .11 M K! Ni . 1 II S.i I I I s

of.i'rct =J at aetorrr-^twjn in (emperwj condition Condition 14GQpn753'CI-t.h»de!cyrr,a-

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

I he elVcel nf tempering ieiv.perature on the retained auMenile content m AI C-77 is also sho'.vn in Fiiiunj VS. The insert m this tigure iiius the results of another investigation which involved a sinelc auMeni/ing temperature but a broader range of tempering temperature-. Increasing tempering temperatures result in decreasing retained austemte contents, due to she precipitation ol larger amounts of carbides and mtermciallic compounds, which raise- the \t. tempera ure. If. prior to tempering, the steel is either ausformed 5(1",, at IOC" 1 (540 Cl or cold worked H>",,. the retained ausiemte contents after 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/vj.

S.4. Gra'.n Size The favourable efleet of decreasing grain size on (he yield stienglh and fracture loughness of high-strength steels prompted the initiation of several studies aimed at controlling the grain size, particularly in AFC-77. Various sequences of therinomechanical trea'ments used to obtain grain refinement in ihi.- sieel are shown in Figure S.y a. Sequence A consists in deforming ihe austenite by at least 70",, in the 2OOO-I5OO F (II00-SI5C) temperature range, cooiing to room temperature, and annealing at 1800 F (980 C). Using the same deformation conditions, efficient grain refinement is obtained when the stee! is reheated for 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 grain after final annealing is much larger (Fig. 8.9 c). The above austenite grain refinement is related to the formation of deformation voids at the interfaces between matrix and hard carbide particles. The voids remain stable up to fairly high temperature (1875 F, i.e. 1025 C for AFC-77) and art as barriers which pin the grain boundaries [8.20]. In AFC-77, the grain size can also be controlled through addition of refining elements such as niobium. With an Nb content of about 0.2%, grain sizes between 4 and 18 {im were obtained for material initially austenized at 1700;F

90 -.1 MM iss \i.\K-\cisr, suns IMIISKU MI I M < I K<,I

(92> C| anu subsequently re-austcni/c

K.5. Ageing Reactions

Ageing reactions in the steels considered here are essentially controlled b\ ;|-,e structure- of the lath martensile obtained after the solution treatment and iv. ihc naiure and content of the alloying elements present. The defect-rich .mien 'structure of il,c manensite governs not only the kinetics of the precipitation reactions (,;. Chapter 2. Section "• > ">. i but also the morphology of the precipitates and even the tvpc .,f pna-c formed. , - b> favouring precipitation of equilibrium rather than meta-Uihle compound- | he mter- metallic phases that can form in stainless maragmg -teel- and allo\- will mm he reMewed with the simple systems serving as an introduction I > the more complex ones upiL.ti of high-strength compositions.

S.5.I. Inm-Chrnmium System

The present general consensus i- that the binary Fe-Cr system exhibits a miscihilny nap at temperatures nelow «33O F (500 Cl. with the b.c.c. solid solution decomposing into nvo isomorphous phases, one rich in chromium u ) and the other rich in iron ui. Figure S.10 shows the latest equilibrium diagram available [X.2I]. which incorporates resulis\if three basic studies [S 22 to K.2-4] concerning the i-phase range. Although the data pertaining to the euiecloid decomposition temperature of this phase and to the boundaries of the various fields are still controversial, the diagram does indicate that steels with chromium contents of 10",, and more can undergo a decomposition reaction on ageiim at temperatures below 930 F (500 C). It has actually been shown [ti.25] that hardening of ihe 17-4 PH steel (l6",,Cr) is due to the formation of a chromium-rich ferrite U i: the same phase 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 relatively short periods at SSO F (475 Cl is sufficient to cause precipitation of the r phase in Fe-Cr alloys containing 20",,Cr. Finally, the structure of an Fe-46Cr alloy aged at V30 I (500 Cl for 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 after ageing at temperatures as low as 660 F(35OC). This would tend to show that decom-

Cr CONTENT, wt.7. 10 .30 40 50 60 70 SO 90 1000

400 Fig. 8.10. 30 40 BO SO 70 Iron-chromium phase diagram. Cr CONTENT, at ".'. After R.P. ELLIOTT [8.21].

91 i i > H \ I I idSi'lMMi 111< . H •>'. V.! N: • 1 11

el IUO'FIMM O-Zhldark-fidd illumination I •/) 1110 F <600"O - 2000b 30,000 Fig. 8.11. — Thin-foil transmission electron micrographs of a 13C'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 intermediate one, the equilibrium product, viz. the c phase, being seldom formed because of its much more 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 between the 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 stable below 1110 F. However, no equilibrium sections of the diagram below HIO'F were presented in the above-mentioned work, in which no account was thus taken of the existence of a decomposition reaction in the Fe-Cr system. In fact, identification of the phases 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 dispersed precipitate forms which is not a Cr-rich ferrite and which, even after 2000 hours, could not 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 identifying accurately 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 the 10",,Co steel (Fig. 8.11 a and h). The fine uniform distribution of the precipitate indicates a homogeneous-type precipitation and reflects the effect of the high dislocation density of the martensite. On ageing at 1110 F (600 C), no precipitate is observed in the l3Cr-10Co alloy; in the 13Cr-20Co alloy, owing to the increase in the degree of supersaturation of the solid solution on account of the higher cobait content, formation of the T phas'j. observed ufter 2000 hours, is preceded by homogeneous precipitation of the "" intermediate " phase (Fig. 8.11 c and

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

TEMPERATURE, °C TFMPERATURE,°C 400 KM 400 600

E 1.0C

osg

08 K

0.7 E — 1920°F(10501'C)-1h,A.C. -1920°F(1050°C)-lh.A.C. -o- DITTO + 930•F(500°C)-tOOh,AC. - DITTO+750°F(400"C)-100h,A.C. 0.6"

100 600 1200 1600 400 BOO 1200 TEMPERATURE, °F TEMPERATURE, "F

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

93 I OH \i l-l O\l MMMi HICiM-SIKl SCiIH SIM IS

HOLDING TIME AT INDICATED TEMPERATURES, hours

1 m M.l 5 1 tftxt i'l" .igeina temperature ,ind tune on hardness of r e-ll>lr-l ••(>' and If-I4(Y UN] slccls. Alter K. BL Ni.AKiir t'i at. [.v/."]. Intual '-arulmon : 1920 I (1050 (_) - 111. A C. are reproduced in Figure X.I3 for the Fe-l9Cr-15Co alloy discussed above, as well as for the Fe-!4Cr-!

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 steels aged for 128 hours at temperatures between 750 and 1110'F (400 and 600 C) as a function

Fig. 8.14.— Hardness of Fe-13Cr-Cc-Mo steels aged for 128 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 System a, A c,A cla

cr-FeCr Tetragonal 8.7995 4.5442 0.5164 (J.-Fe7Mo6 Rhomb.-hex. 4.746 25.78 5.432 FejMo Laves Hexagonal 4.74 7.73 0.613 200 R (Mo-Co-Cr) Rhomb.-hex. 10.903 19.342 1.774 SO 80 BO 120 K0 / (Fe-Cr-Mo) B.c.c. (a-Mn) 8.92 WL'/. Co.wi.% Mo

94 X. S! AINil IK1-. MANAGING SIII.1S IMIVSH Al MMAIILKGY

of the («l."(. Co • \vt.",, Mo) product, ll is seen thai, ut all the ageing iemperatures considered, the hardness increases with increasing (Co • Mo) product. The time to produce peak hardness in such steels decreases as the ageing temperature is raised from 750 to 1200 F (400 to 650 C) [8.34. 8.35]. A considerable number of inlermetallic compounds are likely to form during ageing in lf-Cr-Co-Mo steels depending on their composition and ageing conditions. That this is so has been made clear in a recent review [SJfi] of the equilibrium phases present in the quaiernav> diagram and its associated binary and ternary ssstems. Table K.4 ei'.es the lattice parameters of some of these phases. Table 8.5 summarizes the literature data concerning the phases thought to be responsible for hardening in cobalt-containing stainless maraging steels. As can be seen, a \arieiy of compounds have been quoted. ll should be borne in mind, ho\u"\cr. that because of iheir structural similarity, their positive identification has. in general, onh been possible in overaged specimens [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 are included 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 1 Ccimpositic>n (wt.° ,) Composition Ref. Condition Phases of interiVietallic precipitates (\vt. "„, bal. Fe) i Cr Co Mo Fe

18.31] 13Cr-10Co 930 F (500C) - 2000h intermediate phase _ W.-il) 13Cr-20Co 930 F (500 C)-2000h intermediate phase 1110 F (600 C)-2000h intermediate • a 40.5 15.S I 42.7 l».3l] 13Cr-20Co-2Mo 930'F (500 C) - 2000h intermediate - a 1110 F (600 C)-500h iniermediate • a 34 5 16.8 3.3 43.7 I3Cr-20Co-5Mo IIIO'F (600O- 500h a ^ FeiMo IS-'-'1 Ditto 1110°F (600*C) - 2000h a 4- FeiMo 31.5 17.3 6.9 44.2 [8.38] 13Cr-13Co-4Mo-5Ni 93O°F (500°C) - lOOh clustering —> (Fe.CrhMo 43.1 1.9 28,0 21.5 (5.6NO [8.39] 14Cr-16Co-5 M o-0.5V-0.002C 1200°F (650°C)-2 h 2h Fe^Mo -i- y (traces) 22.1 10.3 29.2 2S.3 (low-carbon AFC-77J 1290DF (700°C)-2 + 2h FeiMo + x (traces) 17.4 !0.1 32.0 2S.2 t400°F (760°C) - 2 + 2h X -r Fe2Mo (traces) 14.7 9.0 34.6 40.7 1500^ (815"C) - 2 4r 2h X + FeiMo (traces) 17.1 9.9 25.5 46.7 [8.39] 14Cr-14L.o-5Mo-0.5V-0.17C 1200°F (650'C) - 2 + 2h MzsCo + FciMo 4- x (traces) (AFC-77) 1290°F (700°C) - 2 4- 2h M23C6 4- Fe^Mo 4- / — — c 1400°F (760 C) - 2 4- 2h M6C + y + M,3C6 15WTF (815=C)-2 4 2h M6C 4- x + M23C6 [8.40] 12Cr-10Co-6Mo-0.1C 1020°F (550' C) - lh M-.C 1110°F (600X) • ih M2C r R i200°F (650'C)- lh M-.C -t- R 1290F (700X)- ih M6C+ R 25.6 1.8 27.8 44.8 W-13] 10Cr-I0Co-6Mo-7Ni 975°F (525°C) - 4h R — — — 1UO°F (600°C)-4h R [i.35] 10Cr-13Co-4.5Mo 93O°F (500°C) - 1200h R 27.6 6.0 31.6 34.9 11IO°F (600'C)- 1200h R 23.2 8.0 27.5 41.4 1290'F (700°C)- 1200h X 24.4 10.0 8.1 57.4 [8.33] 17Cr-8Co-3Mo-8Ni-0.02C cold worked 97% + a - x 84O"F (540°C)- 17h : : : [8.41] 12Cr-14Co-5MooNi-0.5Ti-0.lAl 1560°F (850°C) - lh a + x (grdin boundaries) n [8.6] 12Cr-SCo-2Mo-SNi-0.8Ti-0.02C 880- F (475°C) - lOOOh 7)-Ni3Ti + Fe2Mo 4- x — — — •— I

95 I ;>R\U-l ONT-UNINU HICH-STKhM. I H STKELS

.jl sl.lt> F I'OO (I- IJf'h • 80.000 M 9J0°F • >0O C> - 2000h • 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 temperatures of the order of 1110 F (600 Ci, the precipitates are somewhat richer in molybdenum than the base metal, although not sufficiently so to indicate the formation ol molybdenum imermeiallic compounds such as Fe^Mo or Fe7Mo(,. Ageing at 1200F (650'C) or above gives rise to a very marked enrichment of the precipitates in molybdenum, which substantiates 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 as described 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 the precipitate formed in the 2",,Mo steel after ageing at 930'F (500;C). The precipitate reaction during ageing at this temperature is probably preceded by clustering of solute atoms, mainly Cr and Mo, which is facilitated by the defect-rich microstructure. The clustering mechanism is possibly similar to that found in Fe-Mo alloys [8.42]. Such clustering, followed by the nucleation and growth of Fe2Mo-type particles, has been observed in an Fe-13Cr-13Co-4Mo alloy aged at 930F (500'C) [8.38]. Nucleation of particles on dislocations rather than on martensite lath boundaries seems to be a general feature of precipitation reactions in these systems. In both AFC-77 and a low-carbon version of this steel, significant amounts of the FeiMo Laves 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%Cr steels 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-crystal electron diffraction analysis of samples treated at I200°F (650sC). The composition of this 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-Mo type, but rather of the Fe-Cr-Mo type. It should be noted, however, that the existence of an R phase in the Fe-Cr-Mo ternary system has not been reported [8.44]. The R phase was also observed [8.13, 8.35] in other Cr-Co-Mo steels after ageing at tnode>4lc!y high temperatures (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 AINIFSS V1ARAGING STEF.LS — PHYSICAL METALLURGY

2000 TEMPERATURE, °C 200 400

Q6 OB 10 M 200 600 1000 1400 TEMPERATURE, °F TITANiUM CONTENT, w.t.% l-'ig. 8.1ft. Effect of titanium content on yield and tensile strength Fig. 8.17. — Hardening response of 14Cr-14Co stainless steels as affect- of Fe-12t"r-(tNi-5Co-2Mo-0.02C steel. After K. Bisc.ARDTctal. IX.I7). ed by Mo and C contents. After E.J. Du.is and L. HABRAKIN [t'.S]. Condition: ISfiO F (S5Q C) - 30 mm. W.Q. • 900 F <4H0 Cl-dh.A.C. 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 in AFC-77 on ageing at 1290'F (700;C) or above [8.39]. whereas in Pyromet X-15, the strengthening phase was stated to be R(Co-Cr-Mo) and,or / 18.45]. Cold reduction prior to ageing apparently extends the range of formation of this phase to iower temperatures, 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 by several 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 not included 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 illustrated in 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-Mo steels 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-77 and Pyromet X-12, contain carbon in non-negligible amounts. The data available on these 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 the other two with substantial carbon contents, are compared in Figure 8.17. The carbon-free Cr-Co-Mo steel starts to harden at about 660°F (350°C), and reaches peak hardness at 1020°F (550°C). The molybdenum-free Cr-Co-C steel shows a secondary hardness peak at about 885°F (475CC), which is followed by a rapid decrease in hardness. The steel containing both carbon and molybdenum exhibits a behaviour which is the summation of the two preceding ones. Thus, it is seen that strengthening of Cr-Co-Mo-C steels is mainly 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 carbides may 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" ^*>^^

TEMPERING TEMPERATURE, "F

fie. S.iH. ThiP.-Ioii transn:i*sion electron micrograph I-ig. S.I»*. ! Meet of tempering temperanire on R.T. fracture (nughm1-.*, 0.2",, of ML- after double tempering I- 2 hours) at IOL*> (- V*S arul notched tensile strength ol" AK-77. After R.T. AII.T it til. \S.7\. (?wo Ci. \fter 1. DIIIIKHK M <•/ ill. It.II}. I7.51K) Initial condition: 14(11)1- |7M>Ci-2h • 1W0 F I !(M0 Cl - Ih." () (J. • refrigeration. 2 2h tempers.

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 and 1500 F (650 and 815 C). v.liile MftC cat bides start to occur and finally become predominant in ihe 1400 to 1500 F 1760 to 815 C) temperature range. Figure 8.18 reproduces the microstructure obsened in AFC-77 (the Cr-Co-Mo-C steel in Figure 8.17i after tempering at 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 precipitate as needles at from 1020 to 1200 F (550 to 650 C) [8.40\. There is evidence thai such carbides 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 to proceed 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 associated with secondary carbide precipitation (Fig. 8.17): M^Q, carhides form at 1110 F (600 Cl and above: beyond about 1290 F (700 C). MhC carbides become progressively predominant until soluiioning of all the precipitated phases occurs. This p ittern is similar to 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) temperature range, 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 ageing processes occurring in these steels. In this respect, cobalt plays an important role since this element is capable of producing a martensitic structure in steels with the high chromium contents needed for corrosion resistance. The defect-rich martensite accounts for 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 is observed in Fe-Cr-Co alloys, and is presumably related to the decomposition of the Fe-Cr solid 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 identified with certainty in such alloys, because of its fineness. At higher ageing temperatures, the precipitate coalesces, permitting positive identification of the er phase. The early stages of 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-formation stages of precipitation. For low,molybdenum contents, up to 3.',,, lhe phases formed are of the finely-dispersed i-lype [8.3!. 8.33]. For higher molybdenum contents. \i:.. from 4 to 6",,, this element leads to the formation of other phases : Fe2Mo. R or /. [he latter prevailing at the highest ageing temperatures. The composition of the precipitates tends to confirm the occurrence of Fe^Mo and / phases, or 'H- formation, under specilic conditions, 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 also raise 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 of the matrix, thus increasing the number of dislocations available as sites for precipitation. Finally, cobalt may induce ordering of the matrix through an enrichment effect resulting from 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 the toughness of AFC-77: although some of the numerical data will be used in the next chapter, 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, and structure, on the other. 8.6.1. General The effect of tempering temperature on the yield strength, notched tensile strength and plane-strain fracture toughness of AFC-77 is shown in Figuie 8.19. The notched strength and toughness are excellent up to 700 F (370 C). but both these properties exhibit a sharp drop between 700 and 800 F (370 -and 425C). Other data [8.8] situate the onset of the decrease in fracture toughness at about 800:F. On the other hand, measurements of the mergy absorbed in impact tests on notched specimens show that the rapid loss in ioughness 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 brittle behaviour is characterized by cleavage fracture and not by a predominant intergranular or grain-boundary fracture. Examination of replicas failed to allow identification of the precipitation reactions occurring in the 600-1000°F (3I5-540DC) range; only after tempering = at 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^C^ and M&C (in Figure 8.17, the onset of hardening in the carbon-free steel lies at about 660°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 the decomposition 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 in Figure 8.20, in which Kic has been plotted as a function of tempering temperature, with the retained austenite content as an independent parameter. The beneficial effect of retained austenite is evident. In this respect, it will be recalled that an optimum stable austenite content of 5 to 15% was one of the basic considerations in the development of Pyromet X-23 (see Section 8.3.3).

99 \t T (.UNTUNING HIGH-STRlNCiTH STEELS

(Mi, MN/n' 1S00 IOT1 2O30

50C• " 600 '00 600 900 1000 H0O G TEMPERATURE T

Pig. S ^0. Mltvl i'l relumed aiMi'nife on Iracture I i£. s.;i StroiBih toughncsi rela- loughnev. of AK'-" After I). UIKSIIR |>./wj Iior^hip lor -MC-""* and steels vienv- Condition : au^etr'eJ ai ! "00 to 2(00 h (C5 ed therefrom. After D. UIBSIIB |A'.v|. M 1150 O i.r fig >»!>>. ; 2h temper

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 an austenite area, it may either continue into the austenite or trace a /ig-zag path through interlocking 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 shown that these are in fact arrested on reaching an austenite area, arourd which they branch and propagate 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 to 1800 MN m-) strength range produced by tempering at 500-700 F Cf, ' 370 C); at the higher 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 decreases markedly. So as to overcome this diffic ilty. the composition of the steel was modified by adding 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 the 260.000 psi strength level achieved. This has been attributed not only to the stability of the retained austenite, but also to the occurrence of a stress-induced martensite transformation [S.9]. AFC-260 is another nickel-modified version of AFC-77; however, in this case, the simultaneous decrease in the carbon content from 0.15 to 0.08 "„ tends to negate the 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 function of ultimate tensile strength in Figure 8.21. The steepness of the AFC-77 curve is a reflection of the diminishing austenite content at the higher strength levels. In AFC-260, austenite is 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 retains austenite 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 austenizing treatment, which is desirable to produce significant amounts of retained austenite, reduces the 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- TE 4GFING TEMPERATURE.'C TENSII F STRFHGTH, MM/m2 MB 450 «X) «50 500 _S50_ 1600 T700 1800 1900 1 y-r 1 tr•F(260-C); ; 20 .1 V« - W 80|

700"F(370-C)\ _ V 60* —^i^-jooo'Rsa) C)_ VI 40

soo 900 noo SQO 900 20 240 260 2S0 AGEING TEMPERATURE, "F 5TREN3TH, 103 psi AGEIMG TEMPERATURE, 'F TENSILE

Fig. 8.22. — Strength and toughness of low-carbon AFC-77 Fig. 8.23. — Effect of grain size on strength: tough- as a function of auslenizing treatment and tempering ness relationship for AFC-77. After D. WEBSTER temperature (2 ~ 2h treatment). After D. WEBSTER [8.10]. [8.46]. Fine : soft-tempered - cold rolled; Toughnesses were measured on sub-size specimens coarse : hot rolled; both followed by 1800'F and are thus of relative significance only. (980°C) anneal, refrigeration. 2 -- 2h temper.

lower temperature results in a better compromise between strength and toughness at all tempering temperatures. This is illustrated in Figure 8.22 for sheet of a low-carbon modification of AFC-77 (0.1I5%C); it is seen that the strength level after the two-stage treatment 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 the decreased 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 martensitic steels 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 to obtain grain refinement consisted in soft tempering the material at 1400 F (760C), followed by cold rolling to 50% reduction, and annealing at temperatures between 1800 and 1900T (980 and 1040=C). This produced grain sizes comprised between 2.3 and 60 (Am; the latter dimension is similar to that produced by standard austenizing within the above-mentioned temperature range. Refining the austenite grain from 60 to 2.3 u.m resulted 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 take advantage of the known grain-refining effect of niobium. A 0.2 %Nb addition proved to be an efficient grain refiner; however, grain refinement was practically inoperative when the alloy was treated to ultra-high strength levels (<•/. Fig. 8.21, " Alloy A "). As stated earlier, 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 stainless maraging steels. However, steels which contain fairly large amounts of reactive elements such as titanium (e.g. the Ultrafort grades) require vacuum melting. As is the case for other high-strength steels, consumable-electrode remelting may be of value, but the effects of this melting process are poorly documented. Nevertheless, it is widely held that, for best toughness and ductility in stainless maraging steels (particularly in large sections),

101 lOIUll-ulMMMM. HIGH-STRENGTH STEELS

l \n\ i HOI WORKING OI SIA'NLLSS MARAGING STEELS

\K.'-'' [«./! Heal uniformly at 2050 F. Hot working temperature range is 1700 lo 2050 F (425 to 1120 O. No particular precautions arc necessary on cooling. 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, large forgings and billets should he sen! hot to a furnace and normalized at 1 htxvi "Mi F (S7U-y;5 I I. air cooied to room temperature and then annealed at 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 working sho'.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 necessitates at bast single \acuum melting. The hot-working conditions recommended for several steels of interest here are shown in 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-strength stainless 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 refrigeration treatment is recommended between quenching and ageing. The low-carbon grades do not necessarily 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 Heat treatment Observations Ref.

AFC-77 (annealed) 1400 F (760X1-:-: or 2- 2- 2 h. A.C. For best machinability (Rc 37-40) {9.11 AFC-77 (aged) 1900 F (1040C) - lh. O.Q. / 500°F (260=C) - 2+2h, A.C. For maximum cold formability [9.1] 1900-F (1040 C) - In, O.Q. / —100F (—73 C) - lh / 700°F For service temperatures below <37OC)-2~2h,A.C. 700°F (370°C) » » 1900=F(I040'O- lh.O.Q./ 1100=F (595°C)- 2 + 2h, A.C. For service :emperatures up to 1100T (595°C) AFC-260 (aged) 1900 to 2000'F (1040 to 1095"O, In,O.Q./ — I00=F (—73cC)- For high fracture toughness and 8h / 700 to 800°F (370 to 425"C) - 2+2h, A.C. good stress-corrosion resistance [9.7] » 1900 to 2000"F (1040 to 10955C)- lh,O.Q./--I00°F (—73"O- For maximum room- and high- 8h / 1000 to 1100°F (540 to 595°C) - 2+2h, A.C. temperature strength »

Pyromet X-12 (annealed) 1300°F (705"C) - 2 to 4h, A.C. For best machinability (Rc 30) 19.2] 3 Pyromet X-12 (aged) 1700'F (925'C) - Ih, A.C. / 900 F (480'C) - 2 to 4h, A.C. For maximum strength (Rc 50) [9.2] 1700 F (925 C) - lh, A.C. / 1100F (595 C) - 2 to 4h, A.C. For maximum toughness (Rc 38) »

Pyromet X-l 5 (afe?rt) 1700°F (925°C) - 30min,O.Q. or W.Q.; 1025"F (550X)-4h,A.C. Rc50 19..'] J Pyromet X-23 (aged) 1500 to I800' F(81S to 980"C)-lh, W.O.O.Q. or A.C. / 900 to Rr 45-50 (Rc 29 after sol jtioning) [9.4] 1050 F (480 to 565C) - 2 to 6h, A.C.

Pyromet X-23 (cast) 2100F(lI50'C)-4h. A.C. / 1700F (925°C) -4h. A.C. / 1O25=F — [9.4] (55OC)-4h,A.C. = Llltrafort 401 (aged) 1560 F (85OC) - 30 min, W.Q. / 880"F (475 CJ - 6h, A.C. — [9.5] c I.'itrafort403 (aged) 1560 F (85O°C) - lh, W.Q. / 895"F (480°C) - 6h, A.C. — [9.6]

102 9 STAINLESS MARAGING STEELS - PROCESSING AND PROPERTIES

TABLE O.3. — TYPICAL ROOM-TEMPERATURE MECHANICAL PROPERTIES OF STAiNLESs MAKAUING STEELS

i C.V.N. Hard- Modulus U T.S., 02°, Y.S, Elong., R.inA., impact. A' N.T.S. ness, of elasticity. Steel • Ref. MN A TV J 10-'psi \ in. U.T.S. o,V 10' psi 10' psi "i; ft.lb Re 10" psi 1 II!-1 „•: M ni

AFC-77 255 1758 200 1.179 17 52 18 24 62 68 1.5 50 [9.1. (aged at 700 Ft (AV 3.9) 9.7) AFC-77 290 2000 214 1476 10 32 9 12 23' 25'' 1.0 53 [9.1. (aged at 1100 F) (A', - 3.9) 9.7] AFC-260 224 1544 188 1296 20 56 — 92 101 — [9.7] (aged at 800=F) AFC-260 254 1751 228 1572 14 44 — 61 67 — [9.7] (aged at 1000 F) Pyromet X-12 236 1627 210 1448 13 38 — _ _ — — 48 30 207 (aged at 900 F) Pyromet X-12 182 1255 150 1034 15 50 11 IS — — — 41 [9.2] (aged at HOOF) Pyromet X-15 235 1620 215 1482 17 60 15 20 — 1.4 48 31 214 [9.3] (K,= 10) Pyromet X-23 258 1779 237 1634 15 58 18 24 70 77 1.4 50 29 200 [9.4] (A'(= 10) Ultrafort 401 242 1669 227 1565 !3 60 41" 56'-' 94 103 1.6 — [9.6] (A', = 3.1) Ultrafort 403 245 1689 242 1669 10 60 25» 34» 49-61 54-67 1.5 [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 appropriate modifications of the heat treatment parameters. For example, ageing AFC-77 and AFC-260 at 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-temperature strength. A special feature of Pyroinet X-23 is that, in the fully heat-treated condition, it exhibits a fair amount of retained or reverted austenite distributed uniformly throughout the matrix. The same structural characteristic can be obtained in AFC-77 and the AFC-260 and " Alloy B " grades derived from it; it was shown in Chapter 8 (see Sections 8.2 and 8.3.3) that the use of higher austenizing temperatures or of a double austenizing treatment leads 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 of strength and toughness.

9.2. Properties

9.2.1. Strength/Toughness Relationship Typical room-temperature properties are shown in Table 9.3 for several cobalt-containing stainless 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, for instance, the data for AFC-260).

The few fracture toughness data listed in Table 9.3 indicate that Klc values of 60,000 to 70,000 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 UiiOM -UMNli HKili SlKLNUllI Sill IS

ULTIMATE TENSILE STRENGTH MN/mJ

Tig. **.I, Strength loughfliv relation>hip for \anous hiyh- strcngth steeU.

\H -" ausicni/vd .it 2IH10 i ill After P. WlHSriK |V,V. VV|; due It' Npeeimen configuration, values .iKnc — HXUK)0 psi \ in. !cnj M be opuniiNlic. .-. .: After D. WtnsriH | A "'I. ^XPYROMET I X-23 AFC-7" ausieni/ed at 1900 F ill) n After EJ. n-.Lis ei til. [9~l r. After D. WEBSTER [

AFCOftO (III) After D. WHISTIR 'Alloj !)"(•) (IV) After D. VVfBSTiK 1«./OJ. Ultrafort 401 and L'llrafon 403 after 175 200 225" 250" 275 300 H. BRANUIS and A. VON DIN STEIN>.S ULTIMATE TENSILE r^ENGTH, 103 ps, 1^.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 amounts of austenite retained at the higher ageing temperatures used (975 to 1000 F, i.e., 525 to 54XTC). Ultrafort 401, and AFC-260 aged at a lower temperature, exhibit A'k. values as 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 a drastic drop in toughness in the transverse direction is indicative of uniform austenite distribution. The fracture toughness of various high-strength stainless steels has been plotted in Figure 9.1 as a function of ultimate tensile strength. In addition to the values given in 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-77 austenized 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 the strength increases as the ageing temperature is raised. The shaded area designated II applies 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 austenizing temperature 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 aged at 800 to 1050cF (425 to 565°C) [9.9]; the data given in Table 9.3 [9.7] for this steel are much 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

C AGEING TEMPERATURE,X AGEING TEMPERATURE, C AGEING TEMPERATURE, °C 400 500 C30 /00 150 500 550 2 : 500 525 550 5/5 600 i2Dr—T "T a280

"800 BOO 12O0 1«B 800 900 1000 950 1000 1050 1100 1150 AGEING TEMPERATURE, °F AGEING TEMPERATURE, ° AGEING 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 for Ultrafort401 and Pyromet X-23 are favourably located with respect to the curves just discussed. On the contrary, the fracture toughness and ultimate tensile strength of Pyromet X-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 strength trough observed in AFC-77 aged at 1000cF (540°C) has been attributed to the fact that the ageing reactions which occur at this temperature cause the martensite to contract, thereby setting up internal tensile stresses in the retained austenite regions (whereas, at other ageing temperatures, these internal stresses are compressive) [9.70]. This trough only occurs in steels that contain significant amounts of retained austenite, i.e., those austenized 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 of Ultrafort 401. It has been found that this steel's behaviour is similar following cold deformation 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 sharp one beyond this value (possibly due partly to the development of a fibre texture, though this has not been ascertained): ductility decreases with increasing amounts of reduction but does not drop to prohibitively low levels; retained austenite vanishes after 20% reduction. Ageing of the cold-worked structure at 900°F (480°C) produces a strength increase of the same magnitude as that observed after annealing; material deformed 98% by drawing and then aged exhibits a strength of 355,000 psi (2450 MN/m2).

105 i OH \l I -i UN! \IMM- Mil .11 M Kl M.I II SI I 1 IS

50 40 60 DEFORMATION,". DEFORMATION. %

I i\:. 'O I fleet "f cold work on tensile properties of L'lirafort 401. After k BiM.vkliT jnd W- Sr>Ri |v.^); Condition: 1 5M> h 1X50 Ci - 30 mm. W.Q.; cold deformed b> drjv.mil tdolted lines); aged j! 900 h (4X0 fi 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 the properties of Lltrafort 401 was dealt with in the preceding paragraph. Figure 9.4 shows the relationship between strength and toughness for conventionally treated, strain-aged n and ausformed AFC-77 steel plate [V.o']. Strain ageing involved !0 (, cold work after an initial 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 toughness combinations for final ageing treatments up to 800 F (425 C); ausforming produces a more 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 with intermittent 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 wire cold drawn from 0.05 in. (1.27 mmi initial to 0.0046 in. (0.12 mm) final diameter (92",v reduction), while values of around 500.000 psi (3450 MN.'m2) were still attained in wire with diameters as large as 0.05 in. (1.27 mm), drawn from rod 0.25 in. (6.3 mm) in diameter and aged appropriately [V.I 2. 9.13].

VELD STRENGTH. M GOO 1600 2003 200 200 500'FCSO'C) • 9OO'FK6O-C) 7OO'F07a'C) a IOM"F6WC) aOO'FiaS'C) a nOO"F<595-C) 150 !

, ion 100

V»V. STRAIN A6ED-

CONVENTI0NAL - 1500 ?00 2S0 40 60 80 TO VELD STRENGTH, «P psi REDUCTION IN DIAMETER, V.

Fig. V.4. — Strength/toughness relationship for AFC-77 Fig. 9.5. — LTect of cold reduction in the conventionally heat-treated, ausformed and by drawing combined with intermittent strain-aged conditions. After D. WFBSTER [9.8]. ageing at 900'F (480 C) on tensile All specimens austenized for Ih at 2000°F (1O95X), oil strength of AFC-77 wire. After V.K. quenched, refrigerated at —100'F (—73 C), and tempered for CHANDHOK and A. KASAK [9.12\. 2- 2h at indicated temperature. Ausforming (50% reduction by Initial condition : wire 0.05 rolling) was carried out at 1000'F (540°C) prior to oil quenching. in. (1.27 mm) in diameter aus- Prior to strain ageing (10% reduction by cold rolling), the tenized at 2000°F (1095°C) refrigerated material was tempered for 2h at 50CF (260°C). and tempered at 500°F (260°C).

106 STAINLESS MARAGING STEALS - PROCESSING AND PROPERTIES

TEST TEMPERATURE, "C TEST TEMPERATURE. °C 100 .200 300 400 _0 150 300 iSD 600 -J5Q 0 I5C 300 450 600

-200 200 60) tOO -200 0 200 600 1000 TEST TEMPERATURE, *F b) Tensile and impact properties of Ultrafort 401 and Uitrafort 403. After K. BUNGARDT md W. SPYRA [9.5] and H. BRANDIS and A. VON DEN STEINEN \9.6\, respectively.

TEST TEMPERATURE^ B0 200 an 400 500 600

100 500 700 1B0 TEST TEMPERATURE, °F i Ultimate tensile strength vs. test temperature. Heat treatment as in Table 9.2 (AFC-77 and AFC-260 aged at 1100 and 1000DF. respectively). O AFC-77 AFC-260 After EJ. Duus and ,-IOOOS Mod. 12%Cr L. HASRAKEN [9.75] PH steels 200 W 600 800 BOO 1200 K00 PyrometX-15. After [9.3] TEST TEMPERATURE, °F Pyromet X-23. After [9.4] c) Ultimate tensile strength of AFC-77 wire : Ultrafort 401 After H. BRANDIS and drawn and aged to 550.000 psi 13800 MN/m ). Ultraforl 403 A. VON DEN STEINEN [9.6]. 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 in this chapter is illustrated in Figure 9.6 a. Their temperature capability is generally superior to that of the best cobalt-free precipitation-hardening steels. The temperature dependence of the mechanical properties of Ultrafort 401 and Ultrafort 403 are shown in greater detail in Figure 9.6 h. Both steels are seen to become brittle at — 321 F (—196 C); the rapid decrease 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 the temperature dependence of the AFC-77 cold-worked wire referred to in the preceding section. It is seen that the wire retains considerable strength up to 1000T (540 C): this additional 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 HC

Fig. 9.?.— 100-hour creep-rupture strength of AFC- ?". AFCOM). PyromeiX-15. UhraforUOl and several cobalt-free precipitation-hardening stainless steels. After E.J. Dixis ei al. [9.7]; data point 800 900 «H1 for Ulirafort 401 after H. BRANDIS and A. VON DIN 'EST 'EWPES4TLIRE. T STHMN [9.6]; data for Pyromel X-15 after rVJ).

The 100-hour creep-rupture strength of AFC-77, AFC-260, PyrometX-i5 and several cobalt-free precipitation-hardened steels is shown in Figure 9.7. AFC-77 is superior in this respect to AFC-260 and Pyromet X-15, which are distinctly better than the cobalt-free steels. According to the single data point available for Ultrafort 401, this steel is more or less comparable to the PH steels. Typical creep and stress-rupture data for Pyromet X-15 are given in Figure 9.8. The favourable creep and stress-rupture properties of the various steels just considered indicate that their thermal stability, or ability to retain adequate properties at high temperature, is good. However, most stainless maraging steels will show some thermal instability when 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 strengths and decreased ductility and toughness. In this respect, a fair amount of work has been devoted to AFC-77 in relation to its potential application in supersonic aircraft. Figure 9.9 shows that holding the steel at 900:F (480C) increases its strength moderately. The accompanying decrease in toughness is not thought to be related to this increase but to be due mainly to the removal of the last traces of retained austenite. No variation of properties was observed in the steel after prolonged exposure at 600" F (315 C). Table 9.4

p=T(ZWogt).I0-3fK.h] 17 320

X 32 34 DO 2B0 Xtl 500 EXPOSURE TIME hours p=T(20*tog 0.10-3 [Ra, hi

Fig. 9.9. — Effect of long-time exposure at 900"F Fig. 9.8. — Creep and stress-rupture data for (480°C) on the tensile properties and fracture Pyromet X-15 (Lar on-Mil!er plot). After [9.3\. toughness 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. DLUS el al. [0.7]

As-aged condition After !000-h exposure at 650 F (345"C) Test under 40,000 psi (275 MN/m-) temperature U. r.s., 0.2% Y.S., Elong. U. T.S., 0.2"; Y.S., Elong.. F C 10'psi MN/nfi 10-' psi WAV*.* 10' psi MSInP- 103 psi MS, m- °'o

A) Mi cold reduction prior to ageing

-110 260 1792 208 1434 11 261 1855 225 155i 7 72 22 237 1634 182 1255 11 215 1482 202 1393 10 650 340 234 146 1007 12 242 1669 162 //.'/ 10 B) 10° , cold reduction prior to ageing — 110 —79 294 2027 284 1958 6 72 22 264 1820 255 1758 5 276 1903 268 1848 h 650 340 25. 1731 218 1503 9 262 1806 230 1586 9

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, or cold 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 and 650°F [9.7]. The data confirm the pronounced strengthening effect of a 10% cold reduction prior to ageing, and show that long-time exposure at temperature under stress results in appreciable 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 401 is 10.>,000 psi (735 MN/m2); this corresponds to a fatigue limit to ultimate tensile strength ratio of 0.43 [9.5], in good agreement with the value of 0.45 typically given as valid for high-sirength steels. It has been suggested that the fatigue limit of Ultrafort 401 can be improved, a? is the case for non-stainless maraging steels, by decreasing the solutioning temperature or by applying surface treatments such as nitriding or shot peening. The data listed in Table 9.5 indicate that the fatigue strength of Ultrafort 403 is higher than that of Ultrafort 401 [9.6]. The fatigue strength of AFC-77 bolts was recently shown to be greater than that of the hot-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 403 After H. BRANDIS and A. VON DEN STEINEN [9.6]

7 CTmin omax for rupture in 10 cycles Steel Specimen Omni 10-' psi

Ultrafort 401 smooi1! 0 99 685 » 1 74 510 notched (Kt = 4.1) 0 43 295 » » —1 28 195 Ultrafort 403 notched (Kt — 4.O 0 47 325 J 33 225

109 mini liUMAIMNO HlUHSIRCNUtH STEELS

i ••• o Ml 1 ;Uvi ol ageing temperature on :hc fatigue behaviour,.( \! C-" bolls. TEMPERING TCMPERATURCC After IV WIBSIIR [».!•>]. a $Q ^f ^P -^ - vondr.ion : :iIW I III50 O - Ih. O.Q.: UNO F .f • .liUdi »- Hi.OO.. - HW I I -'VCi- 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) jnJ 4Milhs. rcs?ectiveK

be in excess of the minimum requirements for tvpical aerospace applic/.iins [9.14]. The endurance limit of these bolls was also found to increase with increasing awing temperature (Fig. 4.Kh.

^2 4. Lcrri'swn Resistance The corrosion resistance of stainless maraging steels is structure as well as composition dependent As regards composition, chromium has the greatest imnorianee, and numerous studies on the Fe-Cr system have contributed to establish this fact. The role of cobalt appears 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 the Co 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 the ferritic AISl 434 type steel of similar Cr content [9.11]. Molybdenum has a ?'.rong beneficial effect 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]. The favourable effect of these two elements as well as that of nickel is confirmed by the fact that the resistance of the Fe-19Cr-I5Co steel just referred to in 25°,, H2SO4 can be raised to or above the level of the austenitic AISi 304 steel through simultaneous additions of l.4",,Mo, 1.7",,Cu and 1.3",,Ni. Ageing of stainless maraging steels affect?, their corrosion resistance as a result of the structural changes which they undergo. In particular, precipitation of Cr-rich phases such as R, 7, n. and (r^CrKMo results in decreased corrosion resistance through chromium depletion of the matrix. Generally speaking, the corrosion resistance is observed to decrease with 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 annealed and aged conditions, when exposed for 1000 hours to a 5% salt spray [9.18]; raising the cobalt content to 8 or 15 % did not modify the corrosion resistance of the steel in the annealed 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-2Mo c alloys in 10% H2SO4 was found to be the same; ageing at 930 or 1110°F (500 and 600 O hardly affected the corrosion behaviour of the 10%Co steel, but decreased that of the 20 %Co steel. Cobalt-containing stainless steels have been subjected to a variety of corrosive environments, 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 403 in bar or sheet form exhibits no corrosion after exposure for 3600 hours in artificial sea water {DIN 50 900) or for 78 hours in salt spray (ASTM B 117) [9.6]. The corrosion resistance of Pyromet X-15 in the solution-treated condition is .stated to be at least comparable 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-23 is claimed to be comparable to other precipitation-hardened high-strength steels of similar alloy content [9.20].

9.2.5. Stress-Corrosion Characteristics The stress-corrosion resistance of stainless maraging steels appears to exhibit a complex dependence 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 iisted in Table 9.6. Ageing in the 700-800 F (370-425 C) range results in good stress-corrosion resistance, 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, but are in qualitative agreement with those of Table 9.6. The stress-corrosion resistance passes through a maximum on ageing at 500 F (260 C). and through a minimum on ageing at 900 F (480 C): a marked recovery is observed on ageing at 1200 F (650 C). This variation is identical to that found for th- fracture toughness (see Chapter 8, Fig. 8.19), except that the A'|,. value does not show a recovery on ageing at high temperatures. A similar parallelism is observed between the K\,- and Klscc values : after ageing at 500F, AFC-77 has a KUec value 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-77 under stress involves a hydrogen embrittlement mechanism. The fact that overageing restores the stress-corrosion resistance but not the toughness may be accounted for by assuming that, after ageing at 1200 F, the crack initiation period represents a major part of the total time to failure [9.21]. There is some evidence that retained austenite has a beneficial effect on the stress-corrosion resistance of AFC-77, as it has on its fracture toughness. Actually, fatigue crack growth rates in AFC-77 containing retained austenite were shown to increase only slightly on changing the environment from dry air to 3.5%NaCl solution; in contrast, structures that did not contain retained austenite experienced a sharp increase in crack growth rate when

TABLE 9.6. — STRESS-CORROSION RESISTANCE OF AFC-77 After E.J. DULIS and L. HABRAKEN [9.13] TEMPERING TEMPERATURE, °C 100 2DD 300 a)0 500 I i . i 1 i

Ageing ternperature, Hardness, Time to racture,* 1 days 80 A °F "C Re Specimen 1 Specimen 2 e ' \ | / \ j-eo 700 370 49 >153 >153 o a \ / 750 400 50 >153 >153 y . O FAILURE CAUSED BY PROPASA- \ 1' 111 • • J / \ 800 425 51 .153 TiuN OF THE FATlGLfc CfiftL* ' < / >153 ' 2C — • • FAILED OUTSIDE THE PLANE OF FATIGUE CRACK DUE TO EDGE 900 480 51 4 5 CRACKING I I V 1000 540 52 50 84 n rt 400 BOO BOO 1000 1100 595 52 15 >153 TEMPERING TEMPERATURE, T

Fie. 9.11. — Effect of ageing temperature on siress-corros- * Test conditions : two specimens 0.05 in. (1.27 mm) thick were aged for 2 -t 2 hours at indicated temperatures (initial condition : I900°F (I040°C) - 30 min, O.Q. ion resistance of AFC-77. After R.T. AULT el al. [9.21]. T refrigeration), then loaded as ben| beams to produce a stress of 124.500 psi Centre-notched fatigue-cracked tensile specimens loaded at 90% (860 MN/m-) and exposed in air while subjected to a daily aspersion with a 5% NaO of notch tensile strength ani exposed to acidified 3.5%NaCl 0 solution containing a wetting agent. pOr comparison, cobalt-free stainless steeis were tested under the same conditions: the results were as follows : solution (pH 1.5). ConJition : 1900 F (1040°C) - lh, O.Q.; AM .150 : 14 and 15 days AM 355 : 5 days PH I5-7MO : > 153 days. —100°F (—73°C)-30min; aged for 2+2h at indicated temperatures.

Ill ( OHM r-l/ONTMMNLi HU.ll STRENGTH STUE1.S

I \BI I ')". SI RISS-COKROSION HI SIS! AMI (»: L II KA1OR1 40J IN ALk.-ULU 5',, Nat I. After H. BRA.NDIS and A. VON ots STEINI.N [V.fi]

Applies stress Stress Time Specimen * to fracture, 10'p W.V m- Y.S. days

Te\ts iir r tif/n wrnptrat il'C 1 Smooth 0.90 I ? 242 0.9? .300 Notched (K: 11 0.90 -42 242 0.95 •45 Tests HI :i; F : inn c Smooth 0.70 I4S0 0.90 13, 14

• RL-niovcJ from -1 4 ir.. iW mm> bar.

the same change was made [^.--]. Further siudies have shown that maximum stress-corrosion resistance in AFC-77 is obtained using relatively low austenizing temperatures associated with low ageing temperatures. On the other hand, a 20"„ cold reduction after austenizing at 1900 F (1040 C) has a detrimental effect, even when followed by ageing at 700 to 800'F (370 to 425 O: however, a 10",, cold reduction applied after austenizing at 1800 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 conditions are not identical with those described in Table 9.6, both sets of results are indicative of the good stress-corrosion behaviour of the steels under consideration, at least for the treatments involved. It should be added, however, that a recent evaluation of massive forgings made from AFC-77 and " Alloy B " has not shown their stress-corrosion resistance 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 to exhibit severely degraded flow tolerance in a corrosive environment as the section size is increased, it would appear that the problem of low KUcr values (10,000 to 20,000 psi \ in., i.e., 11 to 22 MNm ' -) experienced with the AFC-77 and "" Alloy B " forgings may be a characteristic of high-strength stainless steels in general.

Fig. 9.12. — Oxidation behaviour of AFC-77 and several cobalt-free high-strength stainless = TIME, hours steels 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 recorded on various steels, including AFC-77, exposed in still air at 1200 F (650 C) are shown in Figure 9.12; the oxidation resistance of AFC-77 appears io be superior to that of the other steels selected for comparison [9.13].

9.3 Secondary Processing

Cold working of stainless managing steels is best performed n the solution-annealed condition 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 in the solution-annealed condition for the low-carbon grades (e.g. Pyromet X-15, Pyromst X-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 such a; Pyromet X-23 can also be machined in the aged condition; however, lower speeds and more rigid work supports should then be used. Surface treatments ich as nitriding increase the wear resistance, and probably also the fatigue 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 hardness of 1340. The nitriding conditions coincide with those for ageing, so that only one treatment is required. Welding, generally in the solution-annealed condition, can be satisfactorily performed by the TIG process with or without a filler metal of matching composition. After welding, a simple ageing treatment is sufficient to achieve a joint strength similar to that of the base metal; however, for optimum strength and ductility (particularly in heavier thicknesses), a post-weld solution treatment is usually recommended. TIG welding of Pyromet X-15 strip 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 of the latter steel, it has also been shown that no susceptibility to hot cracking develops during welding; however, due to low notch toughness in the heat-affected zone after ageing, the joint should be carefully designed [9.24].

9.4. Applications

Cobalt-containing stainless maraging steels have been designed for applications requiring high 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 applications 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 springs operating at elevated temperatures, and extrusion, forging and die-casting dies. Although many of these applications are still under evaluation, it appears that significant a^age may well develop for steels like AFC-77 in the aerospace field as fasteners [9.14] or as structural components with the advent of Mach 3 supersonic aircraft [9.25]. AFC-77 has also been extensively 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 the chemical industry has recently been demonstrated [9.26].

113 (.HUM l-l lAI \i\IM. lilt.II-SIR1 Nt.IH STtflS

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 hardening etiec, cobalt increases the \/< temperature: this factor has been used to advantage in uiHim 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 substructure at higher tempering temperatures, thereby giving rise to a finer precipitation of dislocation-nucleated carbides or intermetallic compounds. Final'.', cobalt may affect the precipitation reactions themselves, e.g. by decreasing the solubility of molybdenum and chromium, and may induce ordering in the matrix. \> a result of these efleets particularly the obtention of a tine dispersion of the strengthening phases, the steels under consideration exhibit attractive combinations of strength and Toughness. The Ai, >'.v. yield strength relationships shown ia Figure 10.1 recapitulate the information presented in Chapters 4, 6 and ''; they clearly confirm the superiority of the 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-Mo steel and the favourable one of the HP 9-4-20 and HP 9-4-30 grades at the 180,000 to

YIELD STRENGTH. MN/m! ,1000 1600 VMO 20C0 2200

160 200 240 2S0 320 CIELD STRENGTH, 10> psi

Fig. 10.1. — Strength/toughness relationship for various categories of high-strength alloys, showing position of cobalt-containing steels. References as in Figures 4.7, 6.1 and 9.1, except for TRIP steels and ,'J-Ti alloys, after A.R. ROSENFIELD and A.J. MCEVILY [10.1].

High values of Klc may not always have been determined according 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 combination of 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 grades also offer good to outstanding compromises. However, the above combinations of strength and toughness are not the only favourable feature of the cobalt-containing high-strength steels. Their other properties lie at least at 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-Mo steels at cryogenic temperatures, and the superior stress-corrosion resistance of the HP 9-4-20 and 10Ni-8Co-Cr-Mo steels. Some properties pertaining to individual steels also deserve a special mention. For instance, the combination of good high-temperature strength and high magnetic induction of the 15Ni-15Co maraging steel is expected to lead to 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 favour their use in tooling and die applications at moderately high temperatures, while the combination of high fatigue strength and stress-corrosion resistance of AFC-77 pleads in favour of its use for high-strength bolts. It has been shown in the preceding chapters that the steels under review do not present any problems as regards processing. In particular, their hot workability, cold formability. machinability and weldability are usually more than adequate. Here again, some specially attractive features of specific steels are worth pointing out. These include the excellent weldability of the HP 9-4-20, 10Ni-8Co-Cr-Mo and I8Ni-Co-Mo grades, and the fact that maraging steel parts can be machined to their final dimensions in the annealed condition, since the subsequent ageing treatment causes minimal dimensional changes and distortion. Appropriate techniques have been developed to fabricate the steels, not only by melting and casting followed by hot and cold working, but also by powder metallurgy methods. Some grades can also be used as castings; this applies, in particular, to a special l8Ni-Co-Mo grade tailored Tor increased fluidity and castability. As regards present and potential uses of cobalt-containing high-strength steels, they are currently restricted to three major fields : (I) aerospace, aeronautic and maiine applications (rocket motor cases, ianding-gear links and other highly-stressed aircraft components, hulls of deep-sea vessels, etc.); (2) structural and machine components (pressure vessels and reactors, bolts and fasteners, bearings, springs); (3) tooling applications (die- casting dies or components thereof, extrusion rams and dies, cold-forming dies, cam followers, tooling fixtures, etc.). On account of their adequate strength retention at moderately high temperatures and their good oxidation and corrosion resistance, the cobalt-containing stainless maraging grades are also being considered for use in steam turbines and jet engines, e.g. as blades and compressor discs. It is difficult to predict the ways in which the cobalt-containing high-strength steels will be used in future, except in the general terms employed in Chapter 1 (see Fig. I.I). However, it can reasonably be expected that the rapid technological advances which characterize our time will make good use of the superior, and sometimes exceptional, combinations of properties possessed by 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 will require the use of the higher-strength maraging grades, and this may eventually lead to very significant consumption of these steels.

115 ( OH \l I-I. ONI A1NING HKIII-STRI NCMII SIEHS

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117 IPIUI I I (IN I \l\INli Illli M KIM. Ill Slit I S

KVIIMI. I' K; ii\. s -\ . .ui!.t I mis. M . \ lo\» Strength jnd \J ID] (.•KOI MM i u. I P . I niiiii n ( I . I IIIII. R I . and SIMONS. I raciutc Ioiie.hm.-v. ol l'N>-4l . Steel JS Allected f> Slum R v . High'SiicniMli Mi-el 'i\i-ii o. DMK Pioicssts and I ci-ipcr ins: .iiul Xii-foniiing. ' •/.*/> /'if <.'"•"' 7 '•*'. *4 Propeities Handbook trtw, I olumbus. ohm

\4 ,'/] Si'ittM. Ci R . l>Aiik<>uski. US and POKIIH. 1 I . Siienglli .'' .'•>; I i'\. I H . .ni.t Rnsfs.su IN. A H . I hi- Propel tics and Mivlo- and loughnev, of le-HINi Alloys Containing l.!ibnn. stiih.li ie ol IONI-21 r-lMo -si o Sled. J Mai,' . ft. "VstlsTli I iMoniium. MoKbdi-iuini and Cobalt. \h I Iiunt. 4. M\\ ll>»"3i .' "I P\,. .-\tK. .1 V Properties o! Pioduc;ion Piodim of HP'M-'O. s Republic S:ecl Research ( entei Puhliiii . Dcicmtvi v Ne i t M siNn. HP. ICIIMIL- Siii-nglh of I ighl C on\'.\ mtio,IIJI ...••-: I KI s. uis. i . VI . I omnbution to the Study of the Heal Steels Between ~l> and 320 I . J Mum . 4. -l\< (I'li.-Ji Itcii'iicni ami Mechanical Piopeities **< s>Ni-4i o Steels. He*. CiK.iv,. J II , I S Sleel Coip . Private t. i U . Sum R. R I . Rim IM k. S Ci . and (>KIiv S J . Rept ! R 12 •" 1'i.iMis I iiw i" Binary Subsiituiional Alloys ol HI C Iron Repl !R I2(li.s-I;. M.IKII II. r'l.h. 1 ikxis ol stijin Rate. Icmperaiuri'and Alloy C ontcnl. hart* [4.IH Pwi.Ufc. J S . PiopeiltfM'l 1'ioduUion Product of lll"'-4-31). 1 ,\ 1/ . h2. i^O i |'»t*»i Republic Steel Re-eauh C en'er Pubhcn . Dvcembci S. |W • _'•': IIK.IIMV.III, 1 |> . [ i si A. A R .and H u i . A M . I he >»Ni-4l o PAM"\I«. 3S. and M\ns. S J. Relationships lleivveen Steel-. HMH Memo 22l>. IJclobcl l'-'M' Mim-ture and Piopeitics in the l>Ni-4l o Alloy System, in ••Striiituie and Piopcities of I Iiiahijih-Sticnglh Steels". •',\ : li"W>v S. I hi Properties of low-Carbon Iion-Nu.ki:l- V l.iommni Maricnsiics. Iran* I I \l t . 2.16. 142^ i I'"- ) \S I M Spec lech Publ No .'"(>. l»v p W.

;.•.'-", i nvNimiik.. V K. HIKIH. J P . and Dins. I J.. I.tlect ot [4.r\ K.\usti. I).. Ki us. S.A., and CIHIIN. VI., l-'km Strength and v. ' ,!l on lenipenng Tool and Alloy Steels. Jiun* ISM. Fracture Toughness oi 9Ni-4Co Steel as AlVected by Strain lempenni! and Ausformmg, Mclat\ lint:. Quart.. 7 (4), 54 (196-). ;.'..'.') vi!AM»u>k. \.K. HtRlH. JR. and Di 1 is, I..!. I Ifect of Cobalt on Carbon Activity and OilfuMvuv in Steel. Iran*. [4I,S\ KAI ISH. [l.. Onus. M.. and Ki uv S A.. Strain ftnipermg of I / \l I. . 124. K5S i Wh2i. Batnite in yNi-4C'ivO.45C" Steel. J. Muter.. 5. KM (19701.

\.'.Z4'. Hoi tiRiMusr. 1 .. and Si HKADIR. H . I he Role of Cobalt in [4. IV\ \'isn\Tvsk^ . C . and SIIIVIIKV>AII), LA, Plane Strain iraclure Carbon steels with Special Reference to Commercial Cobah- Toughness at Room and Sub/em Temperatures, in " Iracuirc C ontaining Alloy Ste^l*. in Particular High-Speed Steels. |r. h. Toughness Testing a( Cryogenic Temperatures ", A.S.T.M Z). Spec. lech. Publ. 49d. i97l, p. 3. 14.-<>] " Mechanical-Property Data on HP 9Ni-M.o-0.20C Steel" 4. (WRBIDK-STRKNGTHENtD STl-.lXS Balletic Memorial insiituie. Columbus, Ohio. September 19"2. PROC KSSING AND PROPtRTlt.S [4.:i] DARSOM. J.M.. IMIIOI. I J . and R. T.P.. I 1st A. A.R.anJ HAIL. A.M.. Fhe9Ni-4Co Propagation in Hiuh Yiclu Streni'th Steels, f.iii'. riimiirf Steels. DMIC Memo 220. October Whf.. Mechanic*. 2, 301 (1971). [4.2\ Putm. II .. Pi«>ir. S.W.. and MAT vs. S.J.. Development and \4. -VI CRIHIISIR. J.W.. and LAMII, I .A., Corrosion latigue Crack Melting of .in Lltrahigh Strength 9",,Ni-4",,Co Steel, in Propagation Studies of Some New Hmh-Strength Structural " t.lectnc Hii'iace Proceedings ". A I. M.I... New York. Vol. 20, Steels. J. Ba*n. tjig.. 91, 570 [W69|. WfO. p. <(l.v [4. 2.1] C ARTHI. C.S.. Crack Kxtension in Several High-Strength Steels [4.J] MAI vs. S.J.. Influence of Impurities am! Related ttlecu, on Loaded in 3.5 "„ Sodium Chloride Solution, Boeing Co. Rept. Strength and Toughness of High Strength Steels. Metals Eng. No. D6-I97^O, November 1967. Qmiri.. 4 (2). 4S (19f.4|. 14.24] STONESIFEU, F.R.. SMITH. H.L., and ROMINL, H.E. Properties [4.4] BIRKLK A.J.. and PORTER. L.F.. The Effect of Cobalt on the of Hot-Pressed lONi-Cr-Mo-Co Steel. Naval Research Labor- Strength and Toughness of Ni-Cr-Mo High Yieid Strength atory Memo Rept. 2065 (AD 699 530), November 1969. Steels. United States Steel Applied Research Laboratory Rept. [4.251 Mt'NOEB, FI.P., High-Strength 9",,Ni-4",,Co Steels and Their on Project No. 40.018-002(16; (AD-600 790), April 1964. Applications. Cahnll. No. 44, 127 (1969). [4.5] RATHBOM . A.M.. Forum : Where We Stand in Welding High- [4.26] Rits. G.D.. and POOLF., S.W.. Welding of Quenched ami Strength Steels. Where Two New Grades Fit into the Picture, Tempered 9Ni-4Co Steels, Welding J.. Res. Suppl., 45- 465s Metal Progr.. 95 (2). 67 {1969). (1966). [4.6] Fe-9Ni-4Co-CrMoV Steels (Code 1221). in -Aerospace Structural Metals Handbook, Vol. I. Ferrous Alloys ". fcds. [4.27] COLTSOLRADIS, D., LAMBERT, N., DRAPIFK, J.M., and HABRAKFN, L., Microstruclure and Properties of 9Ni-^Co V. Weiss and J.G. Sessler. Syracuse University Research Institute. Syracuse. N.Y.. Revised Edition, 1967. High-Strength Steels, Cobalt. No. 37. 192 (1967). [4.28] SAVAS. J.. BECKER, P.C.. and MATAS. S.J., Development of a [4.7] - Properties of Republic HP 9Ni-4Co Steels ". Republic Steel 1X0 ksi Yield Strength Stress-Relievable 9Ni-4Co Filler Metal Corp.. Cleveland. Ohio, !96h; see also Alloy Digests SA-179 Composition. Weldinp J., Rex. Suppl., 48, 479i (1969). (July 19651. SA-27X (July 1972). SA-IK1 (Aug. 1965). SA-270 (Dec. 1971). Engineering Alloys Digest Inc.. Upper Montclair [4.29] UCHIDA, J.M., Welding Evaluation of New High-Strength N.J. Quench-and-Temper Steels, Welding J., Res. Suppl., 45, 31 j [4.H] " Properties of HP 9-4-20 Steel ", Republic Ste.-I Corp. Data (1966). Compilation, September 1972. [4.30] RATHBONE, A.M., Review of Recent Armour Plate Develop- [4.9] STOSESIFER. F.R., and SMITH, H L., Characterization of GTA ments, Blast Furnace SWei Plant, 56, 575 (1968). Weldments in 10Ni-8Co-2Cr-lMo Sleel. Naval Research [4.31] MATAS, S.J., Republic Steel Corp., Private communication, Laboratory Rept. 2466 (AD-746 111), June 1972. 1973.

118 RFFFRENCFS

5. Ni-Co-Mo MAR AGING STKKLS 15.2.'] MIIIAIISI"., JR.. F fleet of Carbon Content on Transformation PHYSICAL MKTAl.lXR(;Y Structure of Iron-22",, Nickel Alloys, in "Fifty Years of Progress in Metallugniphic Techniques". A.ST.M. Spec ['•./] BKAIUIV A J . and (iiiinsi HMIDT, H.J.. An X-Ra> Inveshga- Tech. Puhl. No. 430. I9ISH. p. 250. Imn of the Iron-Rich Nickel-Iron AlUns. J Iruti Steel Imi.. 140, II I I'W I. (.'..V) CiiniiRT. A., and Ovw.v W S.. Diffusionless Transformation in Iron-Nickel. Iron-Chromium and Iron-Silicon Mlovs \5.2] JIIMS. I W'.. and Pi Mni".iv, W I . Free f nergj, and Mctusluhlc Arm Mel.. 10, 45 (19h2i. Slates in the Irun-Niclil and Iron-Manganese Sv^lcms. J Iron Steel ln\t . 163. 121 (I'M1)). [5.25] Hi i/iM,. H.. and KI.IISII-RMASS, J.A.. The Martensitic Tram- formation in Small (0.1-0.3 mmi Iron-Nickel Single Crystals. \5 ?| Itu HI R. ('.(•., Pi ogress with 15",, Nickel StccMor High-Strength Ada Met.. 14. 1'<3 I9 Applications. Metal I'rogr.. 78 (51. 99 (lWtl. [5.21] SWANSON. W.D.. and PARR. J.G.. Transformations in Iron- [5.4] Bn HIR. C.Ci.. 20 and 25",, Nickel Marking Steels, in •• Seminar Nickel Alloys. J. Iron Steel Insl.. 202. 104 (1964). .HI Maraginj: Steels ". International Nickel Co . Inc.. New York. I9h2. p ID [5.27] OVVI.N, W.S.. and WILSOV. FA.. A Note on Mas-,i\e Struclure. i5 .s| HAII . AM . KVMCU of Maraging Sled Developments. Coltulr, in '• Physical Properties of Martensite and Bamiir ". Spec. No. 24. I3K (I'iMi. Rept. No. 93. Iron and Steel Institute. London. 1965. p. 53. ('^l l)n KIK, HI. FASH. J.I.. and GOIDVIAV A.J.. IS",, Nickel [5.2H] BOI R<,K1T. J.. MAITRfflLRRf. P.. M»\>« . J.. and TH(IM»S. B. Muragmg S'cel. lnm\. ISA/.. 55. 5K 114621. Hardening Fe-Ni-Mo and Fe-Ni-Co-Mo Martensitic Alloj- h> Tempering. IRSID Internal P^ept. No. lift. August 1971. (.•>'.."] MAOMI., A . VlAIOlK. I1., DKAIMt.K, J.M.. CiniSlH RAI11S. I)., Part I: Electron Microscope Study of Coherent Precipitation and HABRAKLV L., Microstruciurc. Strength and Toughness in Fe-Ni-Co-Mo Maraging Alloys. Paper presented at of I3NK4O0) Maraging Steel. Cobalt, 1973. 3. 5th International Materials Symposium. Universitv of Califor- |.V#] I uiRitN. S.. The Physical Mctullurgv of Maraging Steels. nia. Berkeley. Calif.. September 13-17. 1971. Met Reviews, 13. 115 |18). [5.2V] MILLER, G.P., and MITCHELL. W.I., Structure of Nickel- [5.v] PKRKAS. M.D.. Structure and Properties of High-Strength Cobalt-Molybdenum Maraging Steels in the Air-Cooled Condi- Maraging Steels. Met. Sci. Heat Treat.. 1970. 558. tion. J. Iron Steel /its/.. 203. 895 (1965). [.'.Ml PI.RKAS. M.D.. and SMTSAR'. V.I.. Eifeet of Alloying Elements [5JO] BONIZEVVSKI. T.. Hydrogen-Induced Delayed Cracking in on the Strengthening of Martensite in Ferro-Nickel Alloys as Maraging Steel. Brit. Weld. J.. 12. 557 (19ft5l. a Result of Heating. Pins. Metals Metallou.. 17 (31. 75 (19641. [5 31] Yf.o. R.B.G., The Effects of Some Alloying Elements on the [5.11] LKiiMiKi. P.. and I_"ASTAI;NI'. J.L.. The Contribution of Elec- Transformation of Fe-22.5°oNi Alloys. Trans. A.I.ME. 111. tron Microscop> to the Slud> of the Technological Propenies 884 (1963). of Maraging Steels. Met. Constr. Mec. 100. 105 (I9M5). [5.32] HAMMOND, CM., The Development of Maragina Stainless [.'./-I Cm Sii. I.-Lis'. and THOMAS. G.. Structure and Properties of Steels Containing Cobalt. Cobalt. No. 25, 195 (19641. Fe-Ni-Co-Ti Maraging Steel. Trans. A.S.M.. 61. 14 (19681. [5.33] MANENC, J.. THIVELLIER. D.. and THOMAS. B.. Hardening of [5.13] C'RIMMISS. P.P.. Evaluation of High-Nickel Maraging Sieels Fe-Ni-Nb Martensitic Alloys by Tempering. IRSID Internal for Application in Large Booster Motor Fabrication, in " Third Rept. No. 116. August 1971. Part 2. Manming Project Review ". Air Force Materials Laboratory [5.3-1] SHYNE, J.C.. ZACKAY. V.F., and SCHMATZ. D.J.. The Strength lech. Doc. Repl. RTD-TDR h.V404K. Nov. !9(-3. p. 9h. of Martensite Formed from Cold-Worked Austenite. Trans. A.S.M., 52, 346 (1960). [.\I4\ SAUOWSKI, t.P.. and DFCKER. R.F , Cast Maraging Steel. AW. Castings. 43 (2). 26 (1963). [5.35] WARD, D.M., MILLER, G.P., and BRIUGES, P.J.. Effect of Martensitic Transformation on Cold Formabilitv of Maraging L5./51 FLOREEN. S.. and DECKER, R.F.. Maraging Steel for 1000 F Steels, Rev. Met., 68, 107 (1971). Service, Trans. A.S.M., 56. 403 (1963). [5..?rJl CHILTON, J.M., and BARTON. C.J., Identification of Strength- 15.16) PETFRS, D.T.. 15"uNi-I5°oCc Muraging Steels Having an ening Precipitates in 18Ni(250) Aluminium, Vanadium aud Improved C ombination of Mechanical and Magnetic Proper- Titanium Maraging Steels, Trans. A.S.M.. 60. 528 (W67). ties, Cobalt. No. 52. 131 (1971). [5.37] MARCUS, H., SCHWARTZ, L.H.. and FINE. M.E., A Study of [5.17] PATTERSON. W.R., and RICHARDSON, L.S.. The Partial Substi- Precipitation in Stainless and Maraging Steels Using the tution of Manganese for Nickel in Maraging Steel. Trans. Mossbauer Effect, Trans. 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119 t OH VI HUM \IMV. ItlliH SIR1 NtiTH STEELS

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V J, The Kinetics and Mecha- metric Analysis of I.o»- temperature Maraging Stainless SteeU, iirtim of the Strengthening of Maraging Steels. AK.VIL lech. Mppon Kinzoku Gakkai-Si, 33. 260 (1969). Rep!. M-3W. Wh5. [5.7J] TtiKLNAiiA, Y. and KINOSHITA, T.. Resistometric Analysis l-s..| OKTIRI. K.. Investigation of Transformation and Precipitation of Low-Temperature Ageing in Maraging Stainless Steels, in l?",,Ni Maraging Steel. Trmt\. ASM.. 59. 2(C . Trans. Japan Insl. Metals, 12, 250 (1971). [5J3] PRIESTER. L., Polentiokinetic Study of Maraging-Type Alloys, [.'..•>'/! BAStRjtF. B.R . and HAI SFR. J.J.. Hardening Mechanisms and Mem. Sci. Rev. Mel., 67. 707 (1970). Delamination Studies of IX"u Nickel Maraue Steels. AFML Tech. Rept. W>-|hh. I9ftd. [5.741 PETERS, D.T., A Study of Austenite Reversion During Ageing of Maraging Steels, Trans. A.S.M., 61, 62 (1968). [.'..'-I MILLFR. G.P.. and MITCHFIL. W.I.. 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121 OH Ml --I ONI UMM, HluliSIRI \(i[H SHITS

"\ \i-(.t>-Mo MAUU.IM. Sll-.l.l.N I*./-'I l.AMUiRt. N.. DiiM'ttu. J.M., ;ind CUITSIH IIADIS. O., Precipi- iY. IXTR\-HU;H SlKIMiTH tiRADKS taimn-llardenini; I obalt-l. opper-MoKhdenum Stainless Steels, ( .•halt. No. 47. M (1970). 1" ;l Ii II*IH. (.• W . and i URNS. R.I 1SNI(35O> Maragmg (.""'./.'I (. AIUN. R I ., and M.WIAK. Ci N . Struclure and Properties of steel, /M>I> I V W . 61. "^>S |l*Si. a Neu High-Strength Maraging Stainless Steel, Cohalt, No. 55. |~ ."} Bi'is. ii. \\ .1 . and C cm \*. I \\ . t lakwnon of a Commercial y; (i-)^;i. 400 ksi lir.tde Mar.igmg Sleel. Paper piescntcd at \ S \I [S 14] lumiMw. G.H . and Hi n . l-.C. ihe 1 Ifeci of Composition Materi.il* 1 ngineerini; Congrc**. IViroil. Oct. 14-1". NivX on the 'temperature of Spontaneous Transformation of Auste- [".>'! Ussiiu. \ \ . Processing of ls",,Ni M.iragina Stc.'l 1350 nue to Mariciisne in IS-M T'\pc Stainless Sleel. Trait.'.. .•t.S.M., Ciradei. -\rm> Malenals and Mechanics Research Cer.ivr 45. 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FISHER, R.M., and SZIRMAE, A., Effect Combination of Strength. Toughness and Stress-Corrosion of 500 C Ageing on the Deformation Behaviour of an Iron- Resistance. Trans. A.S.M., 61. 75 (1968). Chromium Alloy, Trans. A.f.M.E., 230, 676 (1964). [8.S] DLTIS, E.J.. and HABRAKES, L.. Precipitation-Hardening Stain- 1 [$.30] KOSTER, W., and HOFMANN, G., Equilibrium Phenomena in the less Martensitic Steels. Paper presented at "" 4 ' Congres de la Ternary System Ircn-Coball-Chromium, Arch. Eisenhiitlenn:, Communaute Europeenne du Charbon et de 1'Acier ", Com- 30. 249 (1959). mission des Communautes Europeennes, 1968. [SJ1] COUTSOURADIS, D., DRAPIER. J.M., DIDERRICH. E.. and HABRA- [«.P] WEBSTER, D., Development of a High-Strength Stainless Steel KEN. L.. Precipitation Hardening of High-Strength Stainless with Improved Toughness and Ductility. Met. Trcts.. 7. 2097 Steels. Cobalt. No. 36, 144 (1967). (1971). [8.32] Cot—sctiRADis, D., DRAPIER, J.M., DIDERRICH, E., and HABRA- [

122 RLFF.UIiNCF.S

\x..U] DmiKRicn, 1: . DKAIMIR. J.M.. (OITSOI KADIS, 1)., and HAIIBA- [1.8] WIHSTFR, 11.. Increasing the Toughness of the Manensitic KI N. I... Role of CHhall in the Strengthening Mechanisms in the Stainless Steci AFC-77 by Control of Retained Austenite AIC-77 Steel, ,:; ML Tech. Rept. 65-425, I96.S. Content. Ausformmg and Strain Ageing. Trans. A.S.M.. 61, 816(1968). \Kjf] 1 HoMisos, FA., and Wist. O.R.F., Inlermetallic Compound Precipitation in an I•e-IO",,C'r-13",,t'o-5"uMo Alloy, J. Iron [y.9l WIBSTIR. D.. Optimization of Strength and Toughness in Sleet /».»/.. 210, 69| (19721. Two High-Strength Stainless Steels. Met. Trans.,2. 1857(1971). (S.MJ Wisr. O.R.I-'.. The Constitution of Iron-Rich Alloys of the [V.KI] WIIISTF.R, D.. Oe\elopment of a High Strength Stainless Steel with Improved Toughness and Ductility. Met. Trans 2 1097 Fe-Cr-Co-Mo System : A Review. Cobuli, No. 51. 77 (1971;. (I97| IK..)"] Pi ARSON. W.B.. " A Handbook of Lattice Spacings and Struc- [V.ll] Bt.-NtiARDT. K.. SI"»RA. W.. and LKVSAHT/. G.. Investigation tures of Melals and Alloyv. ". Pergumon Piess, London. of Ma;tensitic Ni-Cr and Co-Cr Steels, in •' Steel Strengthening, Vol. I, I95N. and Vol. 2. 1967. Mechanisms "'. Climax Molvbdenum, 7urich. 1969. p. 108. [f.JS] KINOSHITA. T, TciKfNAiiA, Y., and TOVOSIIIMA, T.. Oil the [9./JI CHANUIIOK. V.K.. -\n2, a New High-Strength Maraging Stainless Steel, Cobalt. No. 55. 759 (1969). 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 Optimum 9. STAINLESS MARAGING STEELS — Combination of Strength, Toughness and Stress-Corrosion PROCESSING AND PROPERTIES Resistance, Trans. A.S.M., 61, 75 (1968). [9J] " AFC-77 "", Alloy Digest SS-209 (May 1968), Engineering [9.22] WEBSTER, D.. The Stress-Corrosion Resistance and Fatigui; Alloys Digest Inc., Upper Montclair, N.J. Crack Growth Rate of a High-Strength Martensitic Stainless Steel AFC-77, Met. Trans., 1, 2919 (1970). [9.2] " PyrometX-12 ", Alloy Digest SS-152 (Feb. 1964). Engineering Alloys Digest Inc., Upper Montclair, N.J. [9.23] CATON, R.G.. and CARTER, C.S.. Evaluation of AFC-77 Mar- tensitic Stainless Steel for Airframe Structural Applications, [9.3] " Pyromet X-15 ", Alloy Digest SS-242 (Sept. 1970), Engineer- Air Force Materials Laboratory Tech. Rept. AFML-TR-73-182, ing Alloys Digest Inc., Upper Montclair, N.J.; "Carpenter Sept. 1973. Pyromet X-15", Carpenter Technology Co:p. Technical [9.24] Wu, K.C, and KRINKE. T.A.. Weldability Studies for AFC-77. Note, May 1969. Part 1. Mechanical Properties, Welding J., Res. Stipp!.. [9.4] " Carpenter Pyromet X-23 ", Carpenter Technology Corp. 47, 332i(1968). Preliminary Technical Data Sheet, 1973. [9.25] WEBSTER, D., Optimization of Material Properties to Achieve [9.5] BiJNOARDT. K., and SPYRA, W., Contribution to the Develop- Structure Efficiency Metals Mater., 5, 257 (1971). ment of Stainless Maraging Steels, DEtf-Teeh. Ber., 9, 361 [9.26] SPAHK. H., and KOPP. L.. Possibilities of Application of Very (1969). High Strength Maraging Steels in the Chemical Industry. Chemie-liig'iueur-Tedimk, 41, 1243 (1969). [9.6] BRANDIS, H.. and VON DEN STEINEN. A., D.I'.W., Krefeld, Germ..;iy. Private communications, 1973 and 1974. 10. CONCLUSIONS [9.7] DULIS, H.J., PF.RLMUTTER, I., and KASAK. A.. Heat-Resistanl Fracture-Tough Stainless Steels, WESTEC Pii.rer No. W9-3.3 [10.1] ROSENFIELD, A.R., and MCEVILY, A.J., Some Recent Develop- presented at 1969 Western Metal and Tool Conference and ments in Fatigue and Fracture. Paper presented at 37th Exposition, Los Angeles," Calif., March 10-13, 1969; KASAK. A., Meeting of Structures and Materials Panel, AGARD, Den Private communication, 1973. Haag, Netherlands. 7-12 Oct. 1973.

~~123 I U'MAINIM. HIGH-STRENGTH STEELS~~

~~AUTHOR INDKX~~

v //] 109 [^.5]. 110 [V.5. V./i'l. Ill D -.//]. M3 I'-5]. OAIIKIW-SKI. O.S.. 22 |3..\ 3.6, 3.9), 28 [3.9]. \HRMi-Msos. r.iv. 11. I \t>.->\. ''9 Bi sill. C.C . "I 1|. 42 [4.11). 69 [I5./W]. AM'iM. K.I".. 1'1 I-V.-M DAnosi. F.. 65 [5.7(5]. M-Ki . :: u-' CAIKNS. R.l . 9 [.r.M]. 54 [5.IS). 56 [5./.V|, DA Ci'NHA Bi:i.n. M.. 94 [«..?3]. 95 [8.33\. -\i i r. R.T.. 24 [ 43 [4.14). 4h h'J [ft.//]. " ['./]. 7S[7./i. ^9 [-./]. .SO [7./]. 97 [N.33], 99 [S.33J. .W jj./4). "1 S4 [.v.-j. 9K CAMI'BUL. J.L.. "S |»-""|. DAS, S.K.. 12 [2.39], 15 [2.39\. y4 f.i'.'i. ill W. CAI'I-NIIS. J.M.. M )5.ft/]. 62 [5.ft/|. DAITUVICH, D.P.. 74 |6.29, 6.3(.»], 75 [t>.34], •\1MARH. J.P. >'N CARTiR. i:.S.. 4S [4.23). "5 lrt.3J[. 79 [7.71. 76 [6.29, 6.34). SO ['.7|. 112 [9.23]. DAVIIS. R.G., 14 [2.^5J. B CARI. R.A.. SO [?.!<]. 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). CASTAONF. J.L.. 52 [5. VI. 95 [?t.4!l r BAKER, A.J.. 19 [2.6_\ .\rt.?]. 20 (j.nJ:. 'C\TOS. R.G.. 112 l'J.23). DETERT. K... 20 U.6^], 59 [5.50], 62 , 5.i5.?]. h 58 [5.39]. 59 [5.49], 60 [5.-/V]. l [5.-/9I. CATON. R.L.. 87 [«./.?]. R8 [«./.!). S9 {«./3]. 65 [5.63], 66 [5.50. 5.63]. BASIM. W.R., 58 [5.*)]. 59 [5.401 95 [cV.73], 96 [A',/31. 97 [8.45]. 110 [V./7]. DEVRIES, R.P., 71 [6.15], BVNFRJIE. B.R.. 58 [5.-/5J. 59 (5.5/1 M [.'..'/. Ill [9.20]. DIDERRICH. E., 60 [5.53], 85 [S.//]. 92 [8.31], 5.6/). 62 I5.-/5. 5.5/. 5.(5/]. CHANWOK. V.K.. 17 [2.53]. \S [2.59]. 38 [3.22]. 93 [8.3.'].94 [S.//, S.3.', 8.34], 95 [S.3/, S.3J], 39 [3.22. 3.23], 91 [o..'i5], 95 [«..?W|, 96 [«.39], 96 [8.31], 97 !».//]. 98 [*.//]. 99 [8.31, 8.32), BAKDM. F.J.. 6S [6."\. 97 [».3y]. 98 [.«..f9!, 99 [S.39]. 1.06 [V.i2'\. 110 [9./5]. BARSfiM. J.M.. 4" [4.21). ^4 [ft.j'i. CHAI/DRON. G.. 94 [«.33|, 95 [«..'.»]. 97 [S.J3), DIENTS. G.J.. 15 [2 5/]. BARTON. C.J.. iO [2.30], >R [5.3(5], 60 [5.3(5]. 99 [,?.33]. DlRAN, L.M., 71 [6.14). BASTII-N. P.. 91 [tt.24]. CHENG. I-LIN. 20 [2.65], 52 [5.12]. DRAPIER. J.M.. 23 [3.//]. 24 [3.11], 25 [3.//1, BAYF.R. A.M.. 69 (5.8). 77 [6.8). CHILTON. J.M.. 10 12.30], 58 [5.36], 60 [5.Jtf]. 26 [3.//], 28 13.//]. 29 [3.//]. 30 [3,11). BECKER. P.C.. 49 [-OS]. 49 [4.27], 51 [5.7], 58 [5.44. 5.46], 59 [5.44], CHW-TIAS. J.W.. 7 [2./5], 8 [2.20). BELL. T.. 10 [2.321. 12 [2.32\. 13 [:.ii). 60 [5.53], 61 [5.46J, 62 [5.-W1, 64 [5.7], BHAT. G.K.. 19 [2.6«). CIZEROS, G.. 8 [2..'2], 11 [2.22), 65 [5.77. 5.HO]. 65 [5.7], 11 [7.5]. 78 [7.5]. 79 [7.5], 86-.S./2]. BIF.BFR. C.G.. 50 [5.J. 5.-/]. 54 [5.201. CLAUSING. D.P.. 42 [4.12], 47 [4.12]. 87 [8.12], 92 [«.3/]. 93 [8.32]. 94 [S.32, 8.34], BILBY. B.A.. 8 (2.201- COFFEY. F.J., 75 \fi I'; 95 [5.3/. 8.34]. 96 [S.3/], 97 [8.12], 99 [8.3/, COHEN. J.B.. 64. [5.70]. A'J-'l, 110 [9./5, 9.16). BlRKLt. A.J.. 22 [3.7]. 28 [3.7). 29 [3.7], DIJBOIS, B., 65 [5.78]. 31 [3.7], 37 [3.7], 40 R.4], 42 [4.4]. 69 [(5./0]. COHEN, M.. 8 [2.21]. 27 [.?./•/, 3./5], 28 [i.«I, 29 [3./5I. 30 [3.15], 43 [4.17], 44 [4-/7], DULIS, E.J., 17 [2.53], 18 [2l59], 38 [3.221, BLACKBURN, M.J., 91 [8.271 39 [3.22. 3.23], 84 [5.8], 89 [8.8], 91 [8J6J BOESCH. W.J.. 54 [J./9]. 53 [5.19], 77 [7.21. 45 [4:/ 7, 4.18], 54 [5.22]. CONRAD. H., 20 [2.66). 95 [8.39], 9fi [S.39], 97 [«.«, 8.39], 98 [8.39], 78 [7.2], 7, '..'), 80 [7.2]. 99 [8.8. 8.39], 102 [9.7], 103 [9.7], 104 [9.7], CORN, D.L., 73 [6.24]. 105 [9.7], 106 [9./3], 107 [9./3], 108 [9.7], BONIZEWSKI, T.. 55 [5JO). COLTSOURADIS, D.. 2 [2.2, 2.3), 7 [2.3], 109 [9.7], 111 [9.13], 112 [9.7, 9.13], BONNET, M.. 74 [6.2H]. 9 [2.27. 2.28], 15 [Z27], 16 [2.27], 23 [5.//], 113 [9./3J. BOIJLGER, F.W.. 76 [6.3*]. 24 [3.//]. 25 [3.11], 2fi 13.11], 28 [3.//], 29 [3.//], 30 [3.//], 49 [4.27], 51 15.7], DUNNE, D.P., 7 [2./6J. BOURCEOT. J.. 55 [5.28], 62 [5.28). 63 [5.281, 58 [5.44. 5.46]. 59 [5.44). 60 [5.53], 61 [5.-/(5], DYSON, D.J., 95 [8.40], 96 [8.40], 98 [8.40], 64 [5.28). 62 [5.-A51. 64 [5.7], 65 [5.7], 77 [7.5], 78 [7.5], 99 [8.40]. BRADLEY, A.J.. 50 [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, BRANDIS. H.. 102 [9.6], 103 [9.6). 107 [9.6], S.32, 8.34), 95 [8.31, 8.34], 96 [8.J/], 108 [V.6). 109 [9.(5], 110 [9.6J, 111 [9.6], 97 [8.11, 8.12], 98 [,S.//], 99 [8.31, 8.32], 112 [9,(5j. 110 [9.A5, 9./rt]. EASH, J.T.. 51 [5.6], 60 [5.6]. BRIDGES. P.J., 57 [5J5], 66 [5.35]. COWAN, T.W., 54 [5./9], 58 [5.19]. 11 [7.2], EBELING. D.G.. 96 [8.44). BKIGGS. .I.Z.. 82 [«./], 83 [*./]. 98 [S.I]. 78 [7.2], 79 [7.2], 80 [7.2J. ECONOMOPOULOS, M., 4 [2.9], 5 [2.9], 6 [2.9]. BRisBANt. A.W.. 71 [6.20]. Cox, T.B., 28 [3./i5], 29 [3./6J. EICHELMAN, G.H., 87 [8.14]. BROOK, R.. 69 [6.12]. CRIMMINS, P.P.. 53 [5./3]. ELLIOTT, R.P., 91 [8.21]. Bui, NAM, 65 \5.76). CROOKER, J.W., 47 [4.22]. ELSEA, A.R., 36 [3.20], 40 [4.1], 42 [4.1], BL'STGARDT, K.. 84 [8.6], 86 [8.6]. 87 [8.6], CROWE, D.J., 4 [2.8]. 48 [4./], 49 [4./]. 88 [8.5. 8.17], 89 [S./7], 9i [8.17], 93 [S./7], CUPP, C.R., 6i [5.60], 62 [5.60], 63 [5.60], ENTWJSLE, A.R., 7 [2./9J. 94 [«.<}, 8.17]. 95 [S.6], 97 [S.6, 8.17], 64 [5.60]. 102 [9.5], 105 [?.J, 9.11], 106 [9.5], 107 [9.5, ERICSSON, T., 64 [5.70].

~~124 AUTHOR INDEX~~

HONEYCOMDE. R.W.K.. 4 [2.7]. HORNBOUEN'. E., 18 [2.5*]. 64 [5.68], 96 [8.42]. I--H-LDIN-O, J., 77 [6.41]. HOTZLER, R.K.. 22 [3.10]. 32 [3.10], LACOMBE, P., 8 [2.22], 11 [2.22], 65 [5.80]. FINE, M.f.. 16 [2.55J, 17 13.55]. 58 [J..?7]. LAGNEBORG. R.. 91 [8.28]. 59 [5.37]. 60 [5.37]. 63 (.f.rtrt). 64 [5.ft9]. HOUDREMONT, E , 39 [3.24]. HUIZING, H., 55 [5.25]. LAMBERT. N.. 23 [5.//]. 24 [3.11]. 25 [5.77] FlSHEK, R.M.. 91 [8.29J. 26 [5.77], 28 [3.11]. 29 [5.//]. 30 [3.11], FI.INV. I1 A., 2 |2.4J. HULIT, G.W., 68 [i5.2J. 49 [4.27], 58 [5.46], 61 [5.46]. 62 [5.46] Hi^i-L. F.C., 84 (footnotel. 87 [H.I4]. 86 [8.72], 87 [8.72], 97 [8.12]. 110 [9.16]. 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. LANG. F.H., 76 [6..?s/J. 5.2/|. 58 [5.8|, 60 |5.2/. 5.56]. hi [.?._'/. 5.5VJ. I LANGE, E.A.. 47 ]4.22\. 71 [(5./6J. 1.2 15.X. 5.56. 5.59]. 63 [5.ft. 5.65]. 65 [.57.fl. LEAVERLAND. J.V., 19 [2.61]. (•7 [5.8. 5.31. 5.5V]. 73 [ft.2/]. 74 [6.3V]. IAVNELLI, A.A., 77 [7.J1. 79 [7.3]. LEGENDRE. P., 52 [5.11]. 65 [5.79] 66 K 79] 1? [6.34]. 76 [6.29. 6J4). HI [,S./5], 94 \».I5]. IMHOF. E.J.., 47 [4.21]. 74 [6.25]. 1(11 [«./.?]. 76 [6.57]. IRVINE. K.J., 2 [2./J, 3 [2.6], 4 [2.8. 2.10], LEGER, M.T.. 68 [6.1]. 5 [2.12]. 6 [2.70. 2.12]. 7 [2.72], 15 [2.7, 2.52]. LENNARTZ. G.. 88 [8.77]. 89 [8./71 91 [8 17] 16 [2.7. 2.52), 19 [2/1. 84 [8.4], 88 [8.16]. tiARVvocm. R.D.. 5K [5.43], 59 [5.4;!]. 60 I5.43]. 93 [8.17]. 94 [8.17]. 97 [8.77], 105 [9.11]. 94 [8.16]. 97 [8.16]. 98 [8.16]. CiiiiiuiNS, D.F . 62 [5.62]. 107 [9.7/]. 110 [9.1 i]. Ill [9.77], GILBERT. A.. 55 15..M). LEOPOLD. C.T.. 24 [J./3], J! [4.70], 47 [4.10]. OILI-.WKV, K.P.. 70 [6.78], 71 |6./8]. LESLIE. W.C.. 3 [2.5]. 13 [2..?i5], 14 [2.44]. 16 [2.56], 18 [2J6J, 35 [3.19]. GiuniiERG, A., 67 [J.AV]. JACK, K.H.. 18 [2.57]. LTTTLE, R.T.. 24 [3.73], 41 [4.10]. 47 [4.70]. GOLDMAN. A.J.. 9 [.'.29], 51 [5.(5], 60 [5.6]. JACKSON. J.K., 62 [5.62]. LLEWELLYN, D.T.. 3 [2.6], 15 [232]. 16 [2 51] GOLDSCHMIDT, HJ., 50 [5./J. JANA. S.. 66 [5.82]. 84 [8.4]. GRAAE. A.. 74 [6.271. IECEK. J.. 96 [8.4i]. LOTJERING. G.. 18 [2.58]. ORANGE, R.A., 9 [2.2.?]. JONts. F.W., 50 [5.2]. LUTZ, J.L.. 58 [5.40]. 59 [5.40], GRAY. H.R.. 76 [6.36]. JONES, R.D.. 58 [5.4J], 59 [5.43]. 60 [5.-O]. GREEN, S.J., 35 [i./9]. M GRW.ST.VEUJ, T.P., 24 [3.I3\. 36 13.20], 40 [4.1], 41 [4.70]. 42 [4./]. 47 [4.10]. MCEVILY, A.J,. 114 [70.7]. 4S [,<./], 49 [4.!]. KALISH. D.. 27 [3.14.3.15]. 28 [3.14], 29 [5.7J], MACHUN. E.S.. 13 [2.42J. GROSS, J.H.. 22 [5.4. 3.6]. 31 [3.6]. 37 [3.61. 30 [.?.7J], 43 [4.77]. 44 [4.17]. 45 [4.17. 4.18]. MACIAG. R.J.. 22 [3.70]. 32 [3.70]. 42 [4.13]. 60 [5.57, 5.58], 77 [7.4]. 78 [7.4], 79 [7.4]. GliLBRANSEN, L.B.. ">S [6J51- 80 [7.9]. MCMULUN. J.G.. 96 [8.44]. MAGEE. C.L.. 14 [2.45]. GLRKLIS. J.A.. 76 [6.38]. KARDONSKII. V.M.. 15 [2.50]. KASAK, A.. 17 [2.J.?]. 95 [8.39]. 96 [SJ9]. MAGNEH. A.. 51 [5.7], 58 [5.44], 59 [5.44]. H 97 [8.39]. 98 [8J9], 99 [8.39]. 102 [9.7]. 64 [5.7], 65 [5.7]. 77 [7.5]. 78 [7.5]. 79- [73]. HAURAKEN, I... 2 [2.2. 2.3]. 4 [2.9]. 5 [2.9]. 103 [9.7]. 104 [9.7], 105 [9.7], 106 [9.72], 6 [2.9. 2.I4\, 7 [2J). 9 [2.28]. 23 [3.7 7], 108 [9.7]. 109 [9.7], 112 [9.7]. MAITREPIERRE. P.. 55 [5.28). 62 [5.28]. 24 [3.111 25 [3.//1 26 [3.//J. 28 [3.11], KAUFMAN, L., 8 [2.21]. 54 [5.22]. 63 [5.28]. 64 [5.28]. 29 [J.//], 30 [3.11], 49 [4.27], 51 [.5.71, KELLY, A.. 16 [234], 17 [2.54]. MANENC, J., 55 [5.28]. 56 [5.33]. 62 [5.28], 60 [5.53], 64 [5.7], 65 [5.7] 77 [7.5], 78 [7.5]. 63 [5.28], 64 [5.28]. 79 [7.5]. 84 [S.5], 85 [8.11], 89 [8.8], 92 [8.37], KELLY, P.M., 5 [2.13]. 6 [2.13], II [2.34], 93 [SJ2], 94 [8.77, 8.52, S jfl. 95 [8.31.8.34], 12 [2.34. 2.58]. MANQANELLO, S.J.. 22 [3.6], 31 [3.6], 37 [3.6]. 96 [8.37], 97 [S.S. 8.77], 98 [5.//], 99 [S.8, KENYON, N., 76 [6.39]. MANIAR, G.N., 87 [8.75], 88 [8.75], 89 [8.13], S.3I. 8.32], '06 [9.13], 107 [9.73], 110 [9./5], 95 [8.73], 96 18.73], 111 [9.20]. KEOWN, J.R.. 95 [8.40], 96 [8.40], 98 [8.40], 111 [9.13\. 112 [9.73], 113 [9.13]. MARCINKOWSI, M.J., 91 [8.29]. 99 [8.40]. H,\LL, A.M.. 1 [/./]. 36 [3.20], 40 [4.7]. MARCUS, H., 58 [5.37], 59 [5.37]. 60 [5.37], KESSLER, H., 66 [5.87]. 42 [4.1], 48 [4./], 49 [4.7], 50 [J.5], 77 [6.40]. 63 [5.66], 64 [5.69]. KINOSHITA, T., 64 [5.77, 5.72], 95 [8.38], HAMAKER, l.C, 69 {6.8], 77 [6.8]. MARDER, A.R., 7 [2.77], 10 [2.77], H [2.77], 96 [8.38]. HAMMOND. CM., 9 [2.26], 56 [5.32], 71 [6.75]. 12 [2.17]. KLOSTERMANN, J.A., 55 [J.25J. 84 IS.J], 86 [8.5], 87 [8.5L 110 [9.78, 9.79]. MATAS, S.J., 22 [3.7, 3.2], 35 [3.2]. 36 [3.2]. HASAN, S.Z., 75 [6.35]. KONKOL, P.J., 22 [3.4, 3.5]. 40 [4.2, 4.3], 43 [4.3, 4.16], 45 [4.76], HAUSER, J.J.. 58 [5.45], 59 [5.57]. 61 [5.57, KOPP, L., 113 [9.26]. 46 [4.76], 49 [4.28, 4.31]. 69 [6.9]. 5.61], 62 [5.45, 5.51. 5.61]. KOPPI, W.A., 68 [6.5], 69 [6.5], 75 [6.3]. MAYNOR, H.W., Jr.. 71 [6.17, 6.19]. HAWN. J.M.. 71 (6.20}. KOSTER, W., 92 [8.30]. MELNICK, L.M.. 58 [5.40], 59 [5.40]. HAYNES. A.G., 77 [6.41]. KOUTSKY, J.. 96 [8.43]. MIHALISIN, J.R.. 54 [5.20], 55 [5.23], 60 [5.55], HAYNES, M., 74 [6.28]. KOVESI. P., 19 [2.61]. MILLER, G.P.. 55 [5.29], 57 [5.35]. 58 [5.29], HENRY. R.J., 80 [7.8]. KRAUSS. G.. 7 [2.77], 10 [2.17], !1 [2.77], 59 [5.52], 60 [532], 61 [5.52]. 66 [5.35]. HENTHORNE. M.. 75 [6.33]. 12 [2.17], 66 [5.S5]. MINER, R.E., 62 [5.62]. HIRTII, J.P., 18 [2.59], 38 [3.22], 39 [3.22.3.23], KRINKE, T.A., 113 [9.24]. M..CHELL. W.I., 55 [5.29], 58 [5.29], 59 [532], 91 [8.26] KULA, E.B.. 20 [2.67]. 60 [5.52], 61 [532]. HOIMANN, G., 92 [8.30]. KULIN. S.A., 27 [5.74. i./J], 28 [5.74], MONTUELLE, J . 94 [S.33]. 95 [8.33], 97 [8.33], HOLTMANN, R.B., 84 [8.7], 98'[8.7], 99 [8.7], 29 [3./J], 30 [5.75], 43 [4.77], 44 [4.77], 99 [8.33]. 111 [9.27], 45 [4.77, 4.78]. MOON, D.P., 68 [6.7]. KURDJUMOV, V.G., 14 [2.47], 15 [230]. MOURIKIS, S.. 64 [5.70]. 125 l-OBALT-CONTAlNINli HK.H STRFNGTH STFFl S

\l! V.IK. 111'. :•%,-: '•• H M PK.K n«, R.P.M.. 69 [6.131 74 [6./.?]. ii.A.. 46 [4.19], 47 [4/9). MMRV JR. S4 (S ' PlMi-HRfY. W.!.. 50 [5.2] 73 [6.23] 111 i« :/i Pi ZAK. PP., 71 [A.M|. 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], R *" [5.J9]. \l, IMIM.N. R II . Ir. 12 54|. I" |2.54| RAR('<«-K. SG . 35 {J./vl. SWASSON. W.l). 55 [5.26]. \..\\K. (. J . h'l I'. .'M. "1 I" 'J) RAI-K. H.J.. t.0 [f.57. 5.fX\. 77 [7.4). 78 p.4), S/IRMAI. A.. 91 [,S.29i. Nirnv.. J. n I-'.'"•*I. l-"1 i-'-'J. -'•''!• "9 ['.4), HO (7.9|. 21 |2 !5>*'i. SAVAS. J.. 49 [4-25]. VELTRY, F.J.. 78 [7.6). PATTHISON. VV.R.. 54 [5.17]. ?> [5./T|. ScHMATZ. D.J.. 57 [5.34). VIALATTE, B.. 65 [5.78]. PAI LINA. J P.. 69 [6.W|. SCHRADER, H.. 39 [3.24). VIATOUR. P., 51 [5.7], 58 [5.44,5.46), 59 [5.44], PAXTOS. H.W.. 69 [6.l.i\. 74 (A./.»]. 7? [6..)/), SCHWARTZ, L.H., 58 15.37). 59 [J..?r|, 60 [5.37]. 61 [5.46], 62 [5.46], 64 [5.71, 65 [5.7], 91 [#..V]. 64 [5.69]. 77 [7.5], 78 [7.5], 79 [7.5]. PAVSDN. P.. 9 [:.23\. SERVANT, C, 65 [5.77], VISHNEVSKY, C, 46 [4./9J, 47 [4.19]. 73 [6.2JJ. PEARSON-. W.B.. 5>< 15.47]. 94 |«..^7). SHAPIRO, G . 66 [5.83]. vox DEN STEINFN, A., 102 [9.6]. 103 [9.6]. PEISTRLC. J.N.. M (5/>6). SMMIZU, K., 12 [2.37], 58 [5.41). 107 [9.6], 108 [9.6], 109 [9.6], 110 [9.6], PERKAS. M.D.. 15 [2.50). 52 [5.V. 5.WJ. SHYNE, ).C, 57 [5J4], 111 [9.6], 112 [9.6J. 53 [5.4*J, 58 (5.V]. 59 [5.10]. 67 15.9). SIMONS, R.C., 24 [i./J], 41 [4./0], 47 [4.10]. VOULET, P., 83 [8.2]. PERLMI;TTFR. I.. 102 [9.7]. 103 [9.7). 104 [9.7], SKELTON, M.W., 75 [6.35\. w 105 [9.7], 108 [9.7], 109 [9.7). 112 [9.7], SLUNDER, C.J.. 77 [6.40]. WARD, D.M., 57 [5.J51. 66 [5.35]. PERRY. T.E.. 22 [3.1]. 40 (4.2). SMITH, H.L., 22 [i.«], 29 [3.8], 41 [4.9], WAYMAN, CM., 7 [2.76], 10 [2.JJ], 11 [2.JJ, PETERS. D.T.. 54 [5.16). 61 [5.60], 62 [5.60], 42 [4.9], 48 [4.24], 49 [4.9]. 2.J5). 12 [2.37]. 60 [5.54]. 66 [S.82]. 63 [5.60. 5.<5J. 5.157], 64 [5.60], 65 [5.74], SNAPE, E., 78 [7.6]. WEBSTER, D., 84 [8.10], 85 [8.9,8.10], 89 [8.79], 66 [5.16. 5.741 67 [5.74], 72 [6.22]. SNTTSAR', V.r., 52 [5.10], 59 [5.70]. 90 [8.19, 8.20], 91 [«.9], 100 [8.9, 8.10, 8.19], SOBER, R.J., 35 [3.19]. 101 [8.70,8.46], 104 [9.8,9.9.9.10). 105 [9./0], PETTY. E.R.. 14 [2.43]. 106 [9.8], 108 [9.8], 110 [9.74], 111 [9.22], PICKERING. F.B.. 3 [2.6]. 4 [2.S. _'./0, 2.//], SONNINO, C.B., 75 [6.J5]. 112 [9.22], 113 [9.74, 9J.5]. 6 {2.10. 2.11]. 15 [2.52], 16 [2.52], 84 [S.4]. SPAHN, H., 113 [9.26]. WEST, D.R.F.. 95 [8.35. 8.36], 96 [8.35], PIN'ARD-LEGRY, G.. 5X [5.42). 60 [5.42]. SPEICH, G.R., 10 [2.30, 2.31]. 13 [2.31. 2.36), 97 [S.J5]. 66 [5.42]. 94 [«..?.»]. 95 (»..«]. 97 [SJ.?], 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], WILLIAMS, R.O., 91 [8.22, 8.23]. 99 [#JJJ. 38 [J.9], 39 [J.9], 42 [4.11], 61 [5.59], WILSON, E.A.. 10 [2.32], 12 [2.J2], 13 [2.321 PITLER, R.K., 58 [5.4«]. 60 [5.48). 62 [5.59], 67 [5.59). 55 [5.27]. PITSCH, W.. 66 [5.«/J. SPOONER, S., 60 [5.57]. WOSKOWIAK, A.J., 22 [3.10], 32 [J./O]. POMEY, G.. 91 [8.24]. SPYRA, W., 84 [8.6], 86 [8.6], 87 [8.6], 88 [8.6, Wu, K.C., 113 [9.24]. POMEY. J., 83 [8.2). 8.17], 89 [8.17], 91 [8.17], 93 [8.17], 94 [8.(5, POOLE. S.W., 22 [3.1. 3.3]. 40 [4.2]. 48 [4.26]. 5.77], 95 [8.6], 97 [5.6, S./7], 102 [9.5], PORTER. L.F.. 15 [2.51]. 22 [3.5. 3.6. 3.7. 3.9). 105 [9.5, 9.7/], 106 [9.5], 107 [9.5, 9.7/j, YF.O, R.B.G., 9 [2.24], 56 [5.J/], 57 [5.J7]. 28 [3.7. 3.9), 29 [3.7]. 31 [J.rt. -?.7], 32 [3.9], 109 [9.5], 110 [9.5, ?.//], Ill [9.5, 9.11), 33 [J.9], 34 [i.9], 35 [3.9], 36 [J.9], 37 [3.6, 113 [9.5]. 3.7. 3.9), 38 [J.9], 39 [J.9], 40 [4.4], 42 [4.4, STACEY, A.G., 14 [2.43]. 4.11], 69 [6,70]. STAVROS, AJ., 75 [6.31]. ZACKAY, V.F., 57 [5.34]. PRIESTER, L., 64 [5.73]. S26 SUBJECT INDEX

~~SUBJECT INDEX~~

lfC-77 Haimtic transformations iSre al\o Lower Heat treatment. 40-42. 68. 77-78. 84-85. 102- cold working, 113 bainite. Massive hainile and Upper bain- 103. 113. composition, 82. i««-"«- 4-7- 23-2K. effect on properties and structure. 14-15, 16, heat treatment, 84-85 89-90. 98, 99-101. B. and B, temperatures. 23. 24 17, 18. 27-34. 35, 36. 37-38. 42-43. 71, 102, 103, 105. 110, 111-112. 74-75, 79, 80, 84-85, 89-91. 98, 99-101 hot working, 102. 112. ('«• -V«' AFC-260 105. 109-112. machining, 113. Carbides High-temperature properties. 4ft. 72. 73. 79-80, mechanical properties, 98, 99-101, 103-110. formation. If-. 17. 1X-I9. 25-2". 29-33. 39. 107-109. oxidation resistance. 112, 113. 9?. ^-9H. 99 Hot working, 40. ftK. 77, 102, 112. stress-corrosion resistance. 111-112. morphology. 2ft. 31. 32. 33. 3f. 44. 9K HP 9-4-20 structure, 84,85,89-91,95,96,97,98,99-101. orientation. 5. 1ft. 25-26. 2V,. 30 composition. 22. thermal stability, 108-109. C'arbide-sircngihened steels, general. 22-49. heat treatment. 30, 41, 42-43. thermomechanical treatniem. 90, 101, 106, bainitic transformation. 23-2S. mechanical properties. 26. 27. 30. 42-43, 109, 112. carbide formation. 29-33. 45-48. uses, 108, 109-110, 113. martensilic transformation. 23-24. 2S-29. stress-corrosion resistance, 47-48. welding, 113. strength toughness r.v. siructure relation- structure. 23-24. 25. 28. 29. iFZ-//\ modified (Alloys A and B) ship. 38-39. uses, 49. composition, 82. Carbide-strengthened steels. HP9-4-X grades welding, 48-49. heat treatment, 90-91, 100, 103. ' See also HP 9-4-20. HP 9-1-25. HP 9-4-30 HP 9-4-25 hot working, 112. and HP 9-4-4?) composition, 22, 36. mechanical properties, 100, 101, 104-105. primary processing. 40-42. heat treatment, 40-41. stress-corrosion resistance, 112. secondary processing. 48-49. hot working, 40. structure. 90-91, 100, 101. uses. 49. mechanical properties, 35, 42-45. AFC-26Q (C5Q) Carbide-strengthened steels. Co-modified Iow- structure. 24, 29. composiiion. 82, 100. alloy. 18. 19. thermomechanical treatment, 43-44. heat treatment, 100, 102, 103. 105. Carbide-strengthened steels. Co-modified 5Ni- welding, 48-49. mechanical properties, 100,103-105,107, 108. Cr-Mo HP 9-4-30 structure, 84, 85, 100. heal ireatmem. 29. composition. 22. Alloying elements mechanical properties. 29. 31-32, 37. heat treatment. 30-31. 41, 42-43. aluminium, !5, 52, 56. 84. 88. structure. 31. mechanical properties. 26. 27, 30, 35. 42-43, beryllium, 52. Carbide-strengthened steels. 10Ni-8Co-Cr-Mo 45, 48. carbon. 12, 13, 14, 24, 29. 32, 33, 34, 35. 38. composition. 22. 32. 35. 3ft. stress-corrosion resistance, 47-48. 47, 69, 74, 83, 84. 86, 97. heal treatment, 28. 29. 32-?3. 34. 35. 37-38. structure. 23-26. 29. chromium, 12. 14, 23-24, 34. 36, 38, 39, 52, 41, 42. 43. uses, 49. 53, 56, 81, 84, 86-87. 91-98, 110. mechanical properties. 28, 29. 32, 34, 35. HP 9-4-45 cobalt, 3-4, 7, 9, 12, 14, 15, 16. 18, 19, 23, 37-38, 42, 43. 45, 47. 48. composition, 22. 3b. 29, 31, 33, 34, 37-39, 51, 52, 53, 57, 59, melting. 40. heat treatment, 27, 36, 40-41. 60-64, 65, 67, 81, 83, 8*4, 86-87, 88, 92-99, stress-corrosion resistance, 47-48. hot working, 40. 110, H4. structure, 28. 32-34. mechanical properties, 27-28, 35. 36, 42-48. copper, 15, 52, 84, 87. 97, 110. uses, 49. melting, 42, 43. magnesium, 69. welding, 49. stress-corrosion resistance, 47-48. manganese, 12, 36, 52, 54, 55, 69, 84, 88. Cold working, 48. 76, 78. 113. structure, 24, 27, 30. molybdenum, 15, 18, 23-24, 34, 36, 38, 39, effect on properties and structure, 71-72. 90. thermomechanical treatment, 43^5. 51, 52, 53. 56, 59, 60-64, 65, 67. 84, 86-87, '05, 106. uses, 49. 88, 94-97, 99, 110. Continuous cooling transformation curves, welding, 49. 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- Impurities, 69, 71, 74. 92-94, 100, 110. modified. 15-16. IN-763. 53, 54, 67. niobium, 52, 56, 90, 101. Corrosion resistance, 110-111. Intermetailic compounds, general nitrogen, 18, 84, 86. Creep-rupture strength. 46. 47. 72. 73. 108. formation, 16, 17, 19-20, 57-60, 62-64, 91-97. silicon, 19, 36, 52, 56, 69, 84, 88. morphology, 58. tantalum, 52. D.70, 82. orientation, 58-59, 60, 63. titanium, 52, 53, 56,65,67,75, 84,86,87,97. Intermetailic compounds, types tungsten, 52, 84, 97. Elastic properties, 69, 78. 79, 103. o-FeCr, 91-99. vanadium. 18. 23-24, 36, 52, 56. 84, 88. (FeCrbMo. 95. 1 zirconium, 52. Fatigue strength, 46. 47, 72. 74. 80. 109-110. X(Fe-Cr-Mo) ° -97, 99. AM-367, 82. Ferrite (delta) o-FeMo, 58-uU. Applications. See Uses. effect on properties, 100-101. Fe->Mo, 58-60, 94-97, 99. Austenite reversion, 8, 32-33. 34, 50, 54-55, occurrence, 15, 83-85. U.-FB7MO6, 58-59, 94. 64-67, 88-89. Fracture toughness. See wider Toughness. Fe2(Mo,W), 96. 1, and At temperatures, 8, 9, 24, 56-57, 88. o-FeTi, 58-59. effect on properties, 66, 79-80, 94, 107. Grain size, 62, 85, 90-91. FeVTi, 58. Austenite stabilization. See Retained austenite. effect on properties, 21, 28, 75, 90-91, 101. Ni->FeMo, 59.

~~127 * HIGH STRENGTH STF.FLS~~

\-;.'. P\n:nfi \-l\ Psiomt-I \-:j, Vltral Ii'ii4i)l. I. In,i/,'i I -1(12. and ( Irmfirl 40.1 ,| KI -1 13 '4 N: .\l,. I • "> austenile resersion. NX-S9 N.. I . sv..M-.. -1-. •)' he.11 iriMtnient. ferine (ilellai formanon. X't-X*. 100-101. R.l e-l i-M,>. ••»,.. •« grain size. 90-91. 101 R.M..-1 .'-I • '. '14.'l" '!» m.iilensilic transformation, .S5-X". structure, s'l. t.i-1,4. <•*• |s,iiiie:"-a! fans!,-! iiution times. 24. s4. ss precipitation reactions. 91-99. uses. SO. II s pnm:'r> processing. 101-103 MLiraeme steels. SNI(S(XM i:rade retained auslemte. K9-90 i.om[>oûu»n. ^ \ ^4 .'lei! on precipitation i cut ions. II. Ml. secondary processing. 113 mechanical propeinev. ^ I strength toughness n structure. 99-101. Miiragm • steels. I^Si-'K'o grade ; ;< uses. 1 I 3 'V!|'il,'!,V\ ll'-l 1 . 2V -l- ômpitsilion. ^.*. M Strenglhenuig mechanisms. 2-21. niechanical properties. "2. "' ;T*HVI iii>. I *-i V ^~ hummc phase lraiisliniiiaiions. 4-". 2"-28. >'*'.-:cmpci.iiur;: pn'tvu^v. 2. 2~ precipitation. H>-20. 3s. I,I. ,,-. 94 uses. "; solnl solution. 2-4, IV Hi. 3". 3X. (^). ill, fi7. Maramni; steels. |s\i-l«t'i> i;rade Siv Stress-corrosimi resistance. 4"-4X. '>^. 74-76, K0. 111-112. 1 Marayini: st<-'els. l^NiOMn tudc. \1, M M.^'mnni. . 4.s. ~<\ "v II' 1 Maraging steels t\i-( o-Mei. genetal. -U-M M.11 lensitK transformations .Set i;/\;> I aih m.irtensfte and Iwinned martensiiei. "-lf>. lenipenni: resistance. 15, In. 19. 29. 30-31. -'--4. -S-2''. s(|. <4-s". SS-S". 38-39." W.. \l . and \/i lempenitures. s. 4. 12, 13. Tensile properties 1 preci" ution reactions. 14-20. _"'. I-',. 5h-5~, d4. X4. Sh-S". ductilits. 14. 2h. 2" . 30. 42. 4f>. dft. d8-(S9. 7 7 pnnuir> priVcwiNt:. fiS. ~"--s \lassive baimte 71. 2. S-79. 103. Kl.i. |0h. formation. 4-h strength. 7. 13-15, 19, >. 2". 28-29. 30, 31. 1» strucuire. morpholou\. 4. s 32. 34. 35. 37.39. 42-45. 48. 51, 53. (.1, 66. '-(I. I I- properties, h."", 2>*. 2~. hi, f>8-72. 7K-79. 8S. 97, 98-101. ,'03-10(i. Melting. 4(i. i.S. ~~. 101 -102. 114. elfect on properties. 42. 43. 4". "I. "4. ~\ Thermomechanical treatment. 20-21. lien ircjuncni. hX. "? effect on properties and structure. 43-45, nKvhanical pr(>pcrtics. Ordering, hfi. ^-vlK•t•nst"^ rc'^^lxin fracture. 15, 27. 28. 42. 44-47. 51, 67, 68-72. Maraiiing >icc|v. ISNuM criulc 78-79. 89. 98-101. iOl-W. 114. Powder meiallurg>. M. ^S-"1). impact. 7. 15. 19. 26-3i. 35. ;i7-39, 42. 43, hc.it treatment. <»rv "4-"^. Precipitation reactions See Carbides and fi7, 68-IS9, 78-79. 99. 103. niLvhaniLal prttpcruev 51. hH-f^. ^!-"" Intermeialhc compounds. notch tensile strength yield (ultimate tensilel mcliinu. "I Pynmiei X-12 strength. 35, 3f\ 43. hX-h9, 7R-79, 98, 99, vlros-cittTcsmn rcst^lancc. 7S^6. composition. K2. 103. Murasinu '•led-. ISNiO?0i uruilc heat treatrrenl. 102. Twinned martensite tumpoMiinn. 5.1. hot working. 102. formation, II. 12-13, 55. 85. heat trcaimcnt. hi*. "I. "A-"15 mechanical properties. 103. morphology, II. mechanic;!I propcTtiev 51, f>N-"4. Pyromet X-15 properties. 13-15. niL-luni;. "I. composition. 82. ^!rcss-ci>rro-Hin resistance. "t?-7d. corrosion resistance. 110-111. ÎrucIurc. 5\ 5s. 5(). ft;. ftft. 1,". heal ireatmem. 102, 105. Clirufon 401 iherrmniK-'dunical treutmen:. ~0-~2. hot working. 102. composition. 82. Muraginu steels. IKNifiOOi [;radc machining. 113. corrosion resistance. 110. compiiMlion. '3. ft1*. "4. mechanical properties. 103-105. 107. 108. heal treatment. 102. 109. 113. heal treatment. h.S. "I. 74-"1?. structure. 97. machining, 113. mechanical properties. 51. hN-^4. welding, 113. mechanical properties. 97, 103-109. meltinu. 71. 74. Pynmier X-23 structure. KK, 95. 97. pnvv-der ntctallurg). ^X. 69. composition. 82. thermomechanical treatment. 105-106. sln;ss-C(irro>!iin resistance. 74-7h. corrosion resistance. 111. uses. 113. structure. 5K. fS-i. heat treatment. 102. 103. 105. Ultra),in 402. 82. lhermomechanicai treatmen;. 70-7;. hot working. 102. Vlirafiirr 403 Maraging steels. IKNil.isni grade machining. 113. composition, 82. cold utirking. 7s. mechanical properties. 88. 103-105, 107. corrosion resistance. 110. composition. 53. 54. structure. 8". K8. 95. 97. 99. heat treatment. 102. heat treatment. 77. 79. SO. mechanical properties, 103-104, 107, 109. hot working. 77. stress-corrosion resistance, 112. Retained ausienite, 13. 23-24 27, 32-33. 34, machining. 7x. Upper hainite 65-fi6, 87. 88, 89-90, 105-lOfi. mechanical properties. 51. 7K-XO. formation, 4. 5. effect on properties, fifi. 88, 89, 99-100, morphology, 4-5, 27. melting. 77. 103-104. 105. 111-112. powder metallurgy. 7X-7lJ. properties, 6-7, 27-28. 4J. stress-corrosion resistance, HO. Uses, 49. 72, 77. 80. 113, 115. structure. M). Self-tempering, 28, 29, 39, 41. thermomcchanieul treatment, 79, 80. Stainless maraging steels, general (See also Welding, 48-49. 76, 78, 113. uses. 80. AFC-77. AFC-77, modified (Alloys A and Work hardening, 21. 43-44, 57. 71-72, 76, 79 welding. 78. B). AFC-260, AM-367, D.70, Pyromet 105-106. 113.