Alloying Elements in Steel [2nd ed.]

Table of contents :
Front Cover......Page 1
Title Page (Page iii)......Page 7
Copyright (Page iv)......Page 8
Table of Contents (Page vii)......Page 11
Section 1 (Page 1)......Page 15
Section 2 (Page 13)......Page 27
Section 3 (Page 54)......Page 68
Section 4 (Page 60)......Page 74
Section 5 (Page 68)......Page 82
Section 6 (Page 77)......Page 91
Section 7 (Page 88)......Page 102
Section 8 (Page 113)......Page 127
Section 9 (Page 121)......Page 135
Section 10 (Page 149)......Page 163
Section 11 (Page 176)......Page 190
Section 12 (Page 182)......Page 196
Section 13 (Page 197)......Page 211
Section 14 (Page 224)......Page 238
Section 15 (Page 233)......Page 247
Bibliography (Page 247)......Page 261
Index (Page 248)......Page 262

Citation preview

LIBRARY UNIVERSITY OF CALIFORNIA DAVIS

ALLOYING ELEMENTS

IN STEEL

Bainite

from

of eutectoid austenite at 640 F is martensite formed when the specimen

partial transformation

(350 C). The white background was cooled from 640 F to room temperature at the end of the isothermal, transformation period at 640 F. Steel composition: 0.57% C, 0.82% Mn, 1.16% Ni, 1.07% Cr, and 0.26% Mo. 2500X. An extraordinary example of maximum resolution by light microscopy.

ii

(J. R. Vilella)

ALLOYING ELEMENTS

IN STEEL Second Edition By EDGAR C. BAIN, Vice President (Retired), United States Steel Corporation

and HAROLD W. PAXTON, Associate Professor of Metallurgy, Carnegie Institute of Technology

AMERICAN SOCIETY FOR METALS Metals Park, Ohio Export Sales Distributor:

Reinhold Publishing

LIBRARY

Corporation

UNIVERSITY OF CALIFORNIA DAVIS

Copyright © 1939,

1961 by

AMERICAN SOCIETY FOR METALS Metals Park, Ohio

First Published, 1939 Second Edition, 1961 Second Printing, May 1962

PRINTED IN THE UNITED STATES OF AMERICA

Foreword to the Second Edition When it appeared that the demand would probably continue for "The Alloying Elements in Steel", necessitating a sixth printing, Mr. R. T. Bayless, then Assistant Secretary, American Society for Metals, proposed that the author consider a full review and rewriting of the book. It transpired that the twenty-year-old volume was being used extensively as a textbook by undergraduate college students, and, on this account, the original author concurred with this pro posal on the basis that the book, if continued in print, should ex pound well and accurately the theoretical metallurgy involved. He turned, naturally, to teachers who had used the work as a textbook, for the needed critical assistance in appropriately bringing the book up to date, to overcome such shortcomings as the twenty-year period since its writing had brought to light. Professor H. W. Paxton, at nearby Carnegie Institute in Pittsburgh, having used the book for years, was a logical candidate for such a review, and the original author asked him to collaborate in a re vision. Happily, he accepted the invitation, and the present edition is the result. The current authors hope that any merits in the of view have been original point preserved and that, at the same time, the advances inherent in twenty years of metallurgical progress have been reflected in the changes which have been made in the course of this revision. The illustrations, which tell much of the but they reflect re still of the same as before, are story, basically type cent research in iron-alloy equilibrium diagrams, as well as recent theory on martensite and its tempering. Still emphasized is the basic principle that alloying elements make their main contributions to properties through their influence upon rates of reaction, which in turn influence the most important aspect of steels, that of microstructure. .

.

,

Plttsburgh, August

Edgar C. Bain

1961

v

Foreword to the First Edition

A critical appraisal of the several series of educational lectures pre sented in the past to the members of the American Society for Metals on the occasions of the Annual Metal Congress would undoubtedly emphasize the degree to which the lecturers had succeeded in making admittedly complex technical subjects seem convincingly simple. Nor has this circumstance been the result of either verbal legerdemain or superficiality. It is rather the manifestation of orderly thought, clear writing and, in short, good teaching. With a precedent of this nature well established, the present author could certainly do no less than strive to maintain the tradition; indeed, only to the extent to which this discussion of the roles of the alloying elements in steel adheres to the established standards of simple, logical presentation, may any sat isfaction be taken in its preparation. Most simplifications of comprehensive subjects are accomplished by the device of classification, — the assignment of things and phenomena to categories. Upon the skill in setting up the categories largely de pends the acceptability of the simplification. In the present instance the choice of useful categories may be regarded as more or less obvious but that they be convincing requires a great many exemplifications. In the field of steels carrying alloying elements it is not always easy to secure examples which are not confusing by reason of manifesting more than a single active factor and likewise more than one single effect. In securing an adequate number of examples of the physicalmetallurgical influences exerted by the alloying elements with a single variable exerting a single effect at a time, the author has been greatly assisted by the work of those to whom acknowledgment is made in the footnotes. These acknowledgments cannot, however, adequately express the author's appreciation nor his gratitude. To the staff of the United States Steel Corporation Research Laboratory, and several Research Institutions of the Manufacturing Subsidiary Companies of the United States Steel Corporation, are due special thanks for skill ful experimentation and for data. To Mr. Milton Male, Research Engineer, United States Steel Corporation of Delaware, the author gratefully acknowledges his indebtedness for valued assistance with the drawings illustrating the lectures. Edgar C. Bain

vi

TABLE OF CONTENTS 1

page

(Viewpoint and Scope)

1

Chapter

Introduction

Chapter

Fundamental Characteristics

2

of Steels

4

Allotropy and Binary Alloys of Iron Significance of Equilibrium Diagrams

4 8

Carbon Steels

13

Ferrite-Carbide Aggregate Solid-Solution Austenite Transformation of Austenite Is Controllable Modes of Carbide Dispersion Lamellar Family of Structures Proeutectoid Reaction Acicular Mode of Transformation Nature of Martensite Tempering Spheroidal versus Lamellar Structures Transformation of Austenite at Constant Temperatures Hardenability Grain Size and Hardenability Limitations of Carbon Steel in Hardenability Austenite Grain Size and Final Properties Need for Alloy Steels

14

Chapter

17 18 21

23 27 30 31

34

...

in Alloy Steel

Distribution of the Alloying Elements in Annealed Steel Principles of Distribution General Trends in Distribution of the Elements Effects of Alloying Elements on Ferrite Solid-Solution Hardening Effect in Pure Iron Solid-Solution Hardening in the Presence of Carbide Constant Heat Treatment versus Constant Structure Areas for Investigation Effects of Alloying Elements on Carbide Cementite and the Special Carbides Chromium in Carbides Molybdenum in Carbides vii

43 45 49 51

54

3

Alloying Elements in Unhardened Steel Classification of the Constituents

35 37

57 57 58 58 59 60 61 63 65 66 68 68 71

72

Tungsten in Carbides Vanadium in Carbides

73

Summary

75

Nonmetallic Inclusions and Machinability

75

74

Effects of Alloying Elements in Nonmetallic Inclusions Effects of Alloying Elements in Intermetallic Compounds Nitrides Possibly

a Special

Case

77 77

Effects of Alloying Elements in Dispersed Metallic Particles

Summary Partition of Elements Between Ferrite and Carbide Steels Dominated by Alloy-Bearing Ferrite Low-Alloy High-Strength, Structural Steel Special Low-Carbon High-Alloy Steels Typical Property Changes Induced by Alloying Elements Individuality of the Elements in Unhardened Steel

75

78 78 78 79 79 80 . .

80 82

Chapter 4

Effects of Alloying Elements in Forming Austenite Alloying Elements in Steel Rendered Austenitic for Hardening

88

Importance of Appropriate Heating for Hardening Formation of Austenite by Carbon Diffusion Selecting the Heating Temperature Ternary Space Models and Their Sections Abbreviated Temperature-Constant Alloy Sections Composition Range for Austenite at Constant Temperature Eutectoid Composition and Temperature in Alloy Systems . . Equilibrium versus Rate of Solution (Diffusion) Alloy Distribution in Heated Steels Elements Elements Persistent Elements Usefulness

88

Dissolved in Austenite in Undissolved Carbide Carbides and Grain Size in Nonmetallic Inclusions

of Nonmetallic Inclusions Inclusions and Grain Growth Aluminum-Treated Steels and Grain Growth Summary Chapter 5 Effects of the Elements in Hardening Steel Need for Heat Treatment and Alloys

Characteristics of the Austenite-Martensite Transformation Martensite Formation Retained Austenite Alloys and Martensitic Hardness Carbon and Martensitic Hardness viii

88 89 96 97 100 105 107 110 113 113 113 115 116 116 116 117 121

123 123 123 123 124 125 127

Concepts of Hardenability

128 129

Intensity of Hardening versus Hardenability Measurable Manifestations of Hardenability Estimations and Designations of Hardenability Physics of Hardenability Summary of Hardenability Designations Effects of Alloying Elements on Hardenability Factors Influencing Hardenability Elements in Solution in Austenite Depression of Ar' a Reflection of Hardenability Dissolved Alloy versus Hardenability Elements in Carbides Anomalous Hardening at Low Heating Temperatures Elements in Nonmetallic Inclusions Microscopic Homogeneity and Hardenability Hardenability of Familiar Compositions Factors Extraneous to the Steel Significance of Heating Temperature Unique Structures in Alloy Steels Comparison of Isothermal Transformation Curves Individual Alloy Effect on Bainite Structure Chapter

129 135 138 148 149 150 151

154 156 159 162 164 165 169 171

175 176 177 180

6

Effects of Alloying Elements in Tempering Nature of the Tempering Reaction

182 182 185 188

Time and Temperature Relationships in Tempering Retained Austenite During Tempering Alloying Elements and Softening Retardation of Softening and Secondary Hardness Tempering in the Presence of Carbide-Forming

191

197

Elements

198

Time Interval and Secondary Hardening Secondary Hardness Mechanism Strength at Elevated Temperature Advantages

209 213 221

of Special Dispersions

222

Similarities in Tempered Alloy Steels

222 223

Hardness and Tensile Properties

Loss of Toughness in Intermediate Tempering Low-Temperature Tempering Heat Evolutions in Tempering Loss of Toughness at Higher Tempering Temperatures Tempering of Structures Other Than Martensite Rates of Softening in Three Structures Steels Rendered Inhomogeneous Carburized Steels Hardening Case Carburized Steels Nitrided Steels ix

224

....

224 229 232 233 233 237 237 238 239

Recapitulation

239

Aluminum Chromium Cobalt

243 243 243

Manganese

Molybdenum Nickel

243

244 244

Phosphorus

Silicon Titanium Tungsten Vanadium

244 245 245 246

Index

249

x

Chapter

1

Introduction (Viewpoint and Scope) The story of alloy steel may have its beginning before the dawn of history when primitive man fashioned implements from nickel-bearing meteorites — "iron from Heaven." However, the last forty or fifty years have witnessed more progress in the alloying of special steels and putting them to work effectively than the millenniums preceding. Fascinating as an account of this recent rapid development might be, it must be omitted from this discussion. Also, a chronological survey of alloy steels and their development would, almost certainly, not be the most effective means of becoming acquainted with them. It will be the objective in this book to develop a simple philosophy of alloy steels, for which we shall utilize as many fundamental scientific ob servations as possible. We shall try to build up a concept of the actual function of the various common elements incorporated in alloy steels and to discover how the atoms of these auxiliary elements distribute themselves in the steel, what they accomplish by their presence and how they contrive to alter the properties of the ordinary carbon steels so profoundly. Already it has been necessary to bring in a basis of comparison when the effects of the alloying elements are mentioned. It is natural in considering alloying effects to use as a basis of reference, compari son and evaluation those steels that depend mainly on iron and car bon for their properties and that contain little or no intentional addition of other elements to modify their properties. The mode of approach to the subject will recognize, first, that the plain carbon steels, with their moderate content of manganese and silicon, possess a considerable range of valuable properties which, however, are sub ject to limitations even though elaborate treatments are employed to bring out their potential properties and, second, that these limitations may be extended by the presence of certain other elements, mostly metallic. Interest immediately attaches to the question of which prop erties may be improved and by which elements. Still more funda mental is the question of how these other elements enhance the valuable properties of ordinary carbon steel. In contemplating the property enhancements contributed by alloy ing elements, it is important that measures be employed which are really quantitative in nature. In general, the effects of alloying ele ments are strong when substantial amounts are present; it will be to draw fine distinctions which require hair-splitting ac unnecessary must remain curacy of measurement. Nevertheless, the viewpoint quantitative; otherwise we run the risk of ending up with opinions, conjectures or, at worst, wishful thinking, that reflects a departure

1

ALLOYING ELEMENTS IN STEEL

2

from objectivity

that can be maintained

only when quantitative

are at hand.

data

if

is

by

it

is

it

is

C

is,

a

is

is

it

a

is,

It is essential that the basis of reference (that steels free from or low in special elements) be well understood. This topic has been dis cussed, occasionally at some length, in the books listed at the end of this chapter. It seems appropriate to provide in this book condensed version of the physical metallurgy of iron-carbon alloys which will form an adequate basis for our later discussion of the effects of alloy ing elements. While an inquiry into the alloying elements in steel implies a dis cussion of alloy steels, not regarded as necessary for the present to set forth a purpose finely-drawn definition of alloy steel. The ob of classification of jective knowledge and, subsequent possible, the common alloying elements as to their basic functions in steel — functions that result in contributions to the properties. The effects of the several elements will remain quite the same regardless of what name applied to a steel bearing any of them. As to the carbon steels used as bases of comparison, generally, they will contain less than 1% total of elements other than iron and carbon, little enough to insure relatively unimportant modification of properties. The normal occur rence of manganese may be of detectable significance. In view of the intended scope and space of this book, thoroughly complete survey of all types of alloy steels cannot be provided. Special attention will be directed toward enhancement of mechanical prop erties and to those alloy steels of considerable tonnage which are gen those that show a substantial erally hardened and tempered — that response to heat treatment; this means the steels that carry sufficient carbon and not too much of the alloying element to permit the char acteristic control of transformation of medium and high-carbon steel. Generally speaking, this group of steels contains upwards of about and not more than about 5% alloying element. Considera 0.35% less attention will be devoted to low-carbon steels and to those bly such containing large amounts of alloying element as to preclude the in the ordinary sense, such as Hadfield's man inversions allotropic some of the stainless steels, Invar, and many magnetic ganese steel, alloys. It now well established that unproductive to attempt to cor relate systematically ultimate mechanical properties directly with the without considering the presence of the common alloying elements of the the carbon content and, above all, the proportion element, heat treatment employed and the final structure. Thus, would seem almost misleading to say, without qualification, that any one element contributes, for example, hardness and toughness to steels without stating in what composition and after which treatment. An element does not merely by its presence alone contribute a property, as sugar lends sweetness, without regard for the structure favored the ele ment under specific circumstances.

VIEWPOINT AND SCOPE

3

Accordingly, it seems that a more potent attack would be made on the problem by a consideration of the individual effects of individual elements on structure and thus in turn on properties. The data are not complete and some gaps must remain unfilled, but if the view point is fortunately chosen, the data, as they accumulate, will be found to fit in harmoniously with the plan of inquiry. Most of the discussion will have to do with solubilities, rates of reaction, effective rates of diffusion, and microscopic structures and their persistence. In turn, the relation of structure to useful properties will of necessity be in quired into repeatedly. While it will be possible to consider in detail only those influences of alloying elements mentioned above, the list of properties that may be altered by alloying elements and, under suitable handling, im proved for engineering applications is extensive. Some of the influ ences of alloying elements are as follows:

A. Enhancement of mechanical properties 1. 2. 3.

4. 5. 6. 7.

8. 9.

Increase in strength of steel as manufactured Increase in toughness or plasticity in steel at any minimum hardness or strength Increase of allowable maximum section that may be quench hardened to desired properties Decrease in quench hardening capacity Increase in rate of hardening with cold work Decrease in plasticity at low hardness in the interest of machinability Increase in abrasion resistance or cutting capacity Decrease in warping and cracking in development of desired hardness Improvement of physical properties at either high or low temperatures

B. Enhancement of magnetic properties 1. Increase in initial permeability and maximum induction

Decrease in coercive force, hysteresis and watt loss (magnetically iron) 3. Increase in coercive force and remanence (permanent magnets) 4. Decrease of all magnetic responses 2.

C. Enhancement of chemical

"soft"

inertness

of rusting in moist environment 2. Decrease of attack by air at elevated temperature 3. Decrease of attack by chemical reagents 1 . Decrease

General References Bullens, D. K., "Steel and Its Heat Treatment," John Wiley and Sons, New York, 1948. Sisco, F. T., "Alloys of Iron and Carbon," McGraw-Hill, New York, 1937.

Brick, R. M., and A. loys,"

J Phillips, "The Structure

McGraw-Hill, New York,

of Metals and

Al

1949.

Houdremont, E., "Handbuch der Sonderstahlkunde," lin, 1956.

Springer,

Ber

Chapter

2

Fundamental Characteristics of Steels The element iron, when its impurities are reduced to a few hun per cent, is only slightly harder and stronger than copper. Without too great effort, such iron is producible in a commercial way, but the tonnage represented by such material is relatively small. Its great plasticity and softness are for many applications a disadvantage rather than an advantage, and accordingly nearly all commercial iron is moderately impure, and the processes for winning the iron from the ore and refining it to steel are not designed primarily for the delivery of a highly purified element. In the interest of high yield from the ore, the reducing agent, carbon, is employed plentifully, and later the carbon is eliminated to the desired extent by oxidation, along with dredths

other elements that were reduced from the ore and ald oxidizable. It is this same inexpensive element, carbon, that seemingly influences the properties of iron more profoundly than any other, and generally more favorably. It is allotropy in iron, however, and its attendant effects, particularly with reference to carbon solubility, that make the unparalleled range of properties of steels of the usual possible carbon contents; allotropy is retained in the presence of considerable carbon and other elements, and thereby the capacity for advanta geous heat treatment comes into play. Allotropy and Binary Alloys of Iron. The metal iron, as shown in Fig. 1, exists in two isometric allotropic crystal forms: (1) alpha (a) and delta (8) iron, whose solid solutions are called ferrite (or delta ferrite), and (2) gamma (y) iron, whose solid solution is austenite. The alpha form of pure iron exists below about 1670 F (910 C) and above about 2552 F (1400C), when it is called delta iron. Gamma iron exists at the temperatures between these two ranges. The various alloying elements have widely different solubilities in each of these two forms of iron, a circumstance which brings about unique temperature ranges for the transformations in the presence of added elements. In pure iron, the change from alpha iron to gamma iron occurs at a sin gle temperature as is the case, for example, when a pure substance melts; the presence of any additional elements creates a more or less narrow temperature range over which both forms of iron may exist simultaneously in equilibrium. This circumstance, indicating a parti tion coefficient other than unity, and the fact that this coefficient changes with the concentration of the alloying element give rise to characteristic types of changes in the transformation temperatures of iron alloys; there are two primary types, with subdivisions in these types. All of this is concisely stated in the conventional equilibrium diagram, and for reference the types are illustrated in Fig. 2. How-

4

ALLOYING ELEMENTS IN STEEL

6

ever, the diagrams must be used with great care because in practice they will usually be subject to one or both of the following criticisms: 1.

2.

The steels of commerce contain carbon and other elements which greatly alter the temperatures and compositions involved, except in the case of certain very low-carbon alloys with espe cially high alloy content. The majority of steels, as heated and cooled, actually transform at temperatures far removed from equilibrium temperatures be cause of their very slow rates of reaction near these temperatures. Furthermore, the resulting structures are not predictable from any equilibrium diagram and, as will become clear, the really valuable knowledge concerning alloy systems is that relating to structure.

The diagrams divide the alloying elements into two practical classi fications — austenite stabilizers (Type A) and ferrite stabilizers (Type B). An austenite stabilizer, for example, added to any ironbase system will in general tend to increase the temperature range over which austenite is stable. The characteristics of the transformation in the various binary al loys permit classification as in Fig. 2 and as explained below: Type A, Division I. The alloying element widens the tempera ture range for stable austenite by depressing the alpha-gamma transformation and raising the gamma-delta transformation tem perature. The system shows no iron-rich compounds (or solid solutions in the alloying element) at concentrations encroaching upon the alpha or delta fields. The two- phase (y + 8) region reaches the melting range, and the two-phase (a + y) zone is depressed toward ordinary temperatures. Examples: Mn, Ni, Co. Type A, Division II. Same as Division I above except ironrich compounds (or the solid solution in the alloying element) become stable at compositions encroaching upon the alpha or delta phases. Examples: Cu, Zn, Au, N, and C. Type B, Division I. The alloying element narrows the tempera ture range for stable austenite and finally renders it nonexistent. The austenite zone is completely surrounded by a two-phase by intermetallic (a + y or y + 8) field, that is uninterrupted compounds or solid solutions in the alloying element. Examples: Si, Cr, W, Mo, P, V, Ti, Be, Sn, Sb, As, and Al. Type B, Division II. The same as Division I above except that intermetallic compounds or constituents other than the alpha iron and gamma iron solid solutions make their appearance, in terrupting the "loop" with its enveloping two-phase zone. Ex amples: Ta, Zr, B, S, Ce, and Cb.

FUNDAMENTAL CHARACTERISTICS

Fig. 2. Two possible types, A and B, and the subdivisions, I and of phase equilibrium diagrams for iron alloys. (After Wever)

7

II,

Two of the alloying elements mentioned above have an interesting special effect on steel. Up to 7 or 8% Cr lowers the temperature range of the alpha-gamma transformation; further increments raise it. How ever, the lowering of the gamma-delta transformation temperature in the range up to 8% Cr is more rapid than the lowering of the alphatherefore, to state that gamma temperature, and it is consistent, chromium uniformly narrows the range of austenite stability. Cobalt acts in an opposite way in that it widens the austenite tem perature range even though small concentrations actually raise the alpha-gamma transformation temperature.

ALLOYING ELEMENTS IN STEEL 1600

f\

/

1400

Manganese, 50

40

30

wt

%

60

90

80

70

cf

1534

8Fe-

20

10

1504 C

-

2600

1390 C

SMn 1240 Cx

\ 1245 C

2200

1200

Mn

/

r

yFe

/

/

l

1 1000

Ni 1138C

1

l

11OOC

1 i

1800 ,8Mn

910C £

800

Vw agnetic

769 C

trans) ormatic n

1400

_T27C

/rr 590 C

600 aF e

\ \

400

0

10

Fe

Fig.

3.

20

-H

aMn

\

\

30

40

50 60 Manganese, at. %

70

80

90

iooo

100 Mn

Equilibrium diagram for iron-manganese binary alloys. (Hansen1)

The binary equilibrium diagrams for eight common alloying elements with iron are presented in Fig. 3 to 10, inclusive, f Significance of Equilibrium Diagrams. These diagrams describe the state of affairs at supposed equilibrium in alloy systems. Every known means has been employed to induce all tendencies toward change fThe

binary equilibrium diagrams are used here with the permission of the Book Co., publisher of "Constitution of Binary Alloys", by Max Hansen (Ref. 1 at the end of chapter; superscript numerals refer to lists of references at end of each chapter.)

McGraw-Hill

FUNDAMENTAL CHARACTERISTICS 1600 !

10

20

1

1

Silicon, wt % 30 40

50

60

70

1

1

1

80 90 1

1

1528C FeSi

S1

1400

1430C

1410C

~ Fe5S

k 1381 C

FeSi;

1

1

>oc/ ?00C

1200

r

/ f

4.2

1

1000

I

/

OC

*\

Q.

E

600

l

r

31

a

!

T

U

a' /

A JMagn stic 1 trans orm l

400

l< 1 \

/1 0200

7

S 800

M34

ta

1212 C

j

\l2 10 \ 67'

1 1 1

! r 1 1030 C ! 1 1 11 11 1l 11 1l 11 1 S!25 C ! 1 !

1208

^73.5

Ki

1

^(FeS

c/

69^

2.5

-

) =c

-

— "530C

3t on

200

26

9.5|

l

0

10

1

20

30

40

50 60 Silicon, at. %

70

80

90

100

Fig. 4. Equilibrium diagram for iron-silicon binary alloys. (Hansen1)

ALLOYING ELEMENTS IN STEEL

10

Nickel , wt % 60 50

40

1600

H2600

2200

H 1800

.

e

-\ 1400

1000

600

200

0

10

Fig.

from one

20

30

40

50

Nickel,

Fe 5.

60

70

80

Equilibrium diagram for iron-nickel binary

90

100

Ni

at.% alloys.

(Hansen1)

to another to assert themselves (save perhaps holding the metals at constant temperature for decades or centuries), and all such tendencies have actually been fully realized. In many instances, the difference in temperature required for a reversal of reactions has been narrowed to such an extent that no great error can exist be tween the true equilibrium and the diagram. However, as the inquiry into alloy steels proceeds, it will be apparent how little the student is concerned with equilibrium. Rather than a concern for what tends to state

FUNDAMENTAL CHARACTERISTICS wt %

Chromium, 10

11

20

30

40

50

60

70

80

1

1

1

1

1

1

i

90 11

1830C 1800

^

Cr Fe

y

1600

/

3400

//

✓ ✓

-

3000

-

2600

1539 C 22 , 1507 C 1400

-1390

2200

1200

-13 3

r

a

-19

1000

1800 2

a a) CL E

910C

800 -7.5,

—X

815 C

83 OC

'

\

N

/ \

6O0

p

\\

\

200

Fe

Fig.

6.

20

30

40

1000

H 600

,

\

10

-

*

\

0

1400

\ ?

Magn etic tra nsformc tion^ ( nonequ libriun l)

400

\\

-H

50

60

Chromium,

at. %

\

200

\ 70

Equilibrium diagram for iron-chromium

80

binary

90

100 Cr

alloys.

(Hansen1)

ALLOYING ELEMENTS IN STEEL

12

10 1

1800

50

30 1

1

l

7

/

1

//

1525 C

1400

(390 C/

1

90

1 i

W6Fe7 nr 1

y

1532 C

1

WF e2 W2I "e3 J-

1

20; 1600

Tungsten, wt % 70 80

60

/

-

J

1640C

1

H3000

97.41

1540 C

(W)-j-

/

L 2600

cj

i/

21200 o>

a E

1H2200

2 E

H 1C 40 C

i 1000 910C

800

|

r

78CMa

gnetic

\

j

K

r

i

-i

ii il

i * i

f A\

!

tr 3

J

=.1000

/ / / f

910C -

800

\

✓

840 :

^680 Magnetic

//

\\

\

i

\

\

\

/

.

\

\

cr

-

i

trans ormation

1

I 600

-

i

!

400 i

0 Fe

Fig.

9.

10

20

30

40

50 Vanadium,

60 at. %

Equilibrium diagram for iron-vanadium

70

80

90

100

V

binary alloys. (Hansen1)

tions are altered for producing a specific structure. We might well say that a clear understanding of the behavior of carbon steels is imperative in discussing alloy steels. Ferrite-Carbide Aggregate. Carbon steel, as purchased to be formed

FUNDAMENTAL CHARACTERISTICS 10 1

1800

20

30

1

l

wt %

Columbium,

40

50

60

70

1

i

i

1

CbFe2

,« 5C 7

1

i

J

\\ \ 2 75

1

r

-

2600

1560

-55

1534 C

1400

3000

\ I

i

1600

-

i j

1360C

]

11.5

!

!

i j

122( )C

1

2200

1

°1200 at o. E

J

E

c

1 1

98

1000

1800

I

1

!

!

1

Ka 1 1

800

-

1

1400

|

j

600

0

Fa

Fig.

10.

10

20

j

30 Columbium,

40 o

%

Equilibrium diagram for iron-columbium

50

60 Cb

binary alloys. (Hansen1

400

1

0.2

1

1

1

1

0.4

1

1

10 0.6 0.8 Carbon, %

1.2

—1

1

1

1.4

1.6

Fig. 11. Solubility of carbon in iron as influenced by temperature. Composition is that of commercial steel rather than pure iron.

is

a

is

is

C

a

is,

and shaped and as it usually enters service, is really a mixture of mod erately pure iron with iron carbide, Fe3C (cementite). It may contain some traces of other elements replacing part of the iron. There may also be small particles of nonmetallic substances, inasmuch as steel must be melted and refined in refractories or under slags, the solu bilities of which, small as they are, are greater in steel at high tem peratures than at low. Indeed, it now appears that if there were a chemically pure iron and carbon alloy readily available, it would rarely be employed until a well-advised judicious addition of foreign from the metal element had been made. Commercial carbon steel lurgical standpoint, dispersion of iron carbide in ferrite. The proportion of iron carbide depends very closely on the carbon content, soluble in ferrite at the maximum for only about 0.020% (around 1340 F), and at the lowest temperature at which carbon certainly appreciably less may be rejected from ferrite, the solubility than 0.01%. As working rule, the percentage of cementite in an the per cent carbon content less 0.02% times annealed carbon steel

FUNDAMENTAL CHARACTERISTICS 1550

17

i

99.5% austenite

\

\

\

\

840 Homogeneous austenite 820

\

1500

0.5 % austenite

\

1450

\

\

o.

\

\

Austenite and residual carbide

\\Austenite with carbon

\

\

inhomogeneities

800

\X

u V

780

=

o o

a.

E

E

1400

760

H

H740 1350

H720 1

10

100

1000

10,000

Time, sec rate-temperature curves for commercial plain Fig. 12. Austenitizing carbon eutectoid steel. Prior treatment was normalizing from 1610 F (875 C) ; initial structure, fine pearlite. First curve at left shows be ginning of disappearance of pearlite; second curve, final disappearance of pearlite; third curve, final disappearance of carbide; fourth curve, final disappearance of carbon concentration gradients.2

or subtract 0.3% from the product of the carbon content times 15. Solid-Solution Austenite. Upon heating, the carbon solubility in creases. At the temperature of the formation of austenite in the pres ence of carbon, this increase is abrupt. The solubility of carbon in average carbon steels is shown in Fig. 11. The rate at which the car bon dissolves has nothing to do with this diagram which implies a 15,

close approach to equilibrium, a state that even in carbon steels is often not fully reached in a matter of hours. This rate of solution is shown in Fig. 12 for a typical eutectoid steel austenitized at various

ALLOYING ELEMENTS IN STEEL

18

2800

I

0.20

0.40

0.60

0.80

1.00

1.20

1.40

Carbon, %

Fig. 13. Temperature and composition limits for the formation of pure austenite (shaded ) . Regions for ferrite and heterogeneous con stitution are also shown. The diagram is representative of commercial compositions of carbon steel.

temperatures. To secure a complete conversion of the metallic con stituents of the carbon steels to the single constituent, austenite, re quires a heating into the temperature range that is shaded in Fig. 13. The very minimum temperature range for the solution of the carbide is seldom employed because the time involved may be too long. It is in this wholly austenitic range that much of the forging and rolling of steels is done, but the higher carbon steels (above about 0.90% C) are not always heated in this range for heat treatment because it is often desirable not to have all the iron carbide dissolved. Transformation of Austenite is Controllable. The science of the heat treatment of steel properly begins with the steel at an elevated temperature with all or most of it in the austenitic condition. Here the austenite has two characteristics of interest: (a) composition and

FUNDAMENTAL CHARACTERISTICS 1800 I

1

1

1

1

1

1

1

1

1

19 ,

3000

2600

2200 £

H 1800

A

1400

1000

Fig.

14.

The iron-iron

carbide

diagram.

of composition, and (b) a specific will revert to ferrite and carbide when it

homogeneity austenite

(Hansen1)

grain size. The is cooled and by

is

is,

adjusting the conditions of cooling, specifically the rate, a control of the microscopic structure is possible; that the mode of distribution of the hard carbide constituent in the otherwise soft ferrite under

a

is is

it

is,

it

is

considerable control. One of the rates at which steel may be cooled from the austenitic state the more or less automatically-determined rate at which rolled products will cool when laid out on the cooling bed after the last pass or hot operation. Technically this not a heat because in practice treatment, but metallurgically reproduci ble handling of a given section and results in definite mode of trans-

ALLOYING ELEMENTS IN STEEL

20

0

0.2

0.4

0.6 0.8 Carbon, %

1.0

1.2

Fig. 15. The transformation temperatures in the pure iron-carbon alloys as influenced by heating and cooling at 0.125° C per minute. The probable equilibrium temperatures for the several phases are also shown. (Mehl and Wells3)

is

by

by

is

is

5

is,

formation with resultant property acquisition suited to certain uses. Accordingly, the so-called "as-rolled" condition must be regarded as resulting from one of the heat treatments even though no particular effort is required to achieve the rate of cooling. When carbon steels are cooled at an exceedingly slow rate (a few will occur fairly degrees per hour), the transformation of austenite in accord with within or at most 10° of the that nearly Fig. 13, but more practical when are cooled at equilibrium temperature, they The different. rates, temperature of reaction corresponding quite shown in Fig. 14. diagram for the pure iron and carbon system The lag in the transformations — both on cooling and heating — in Mehl and Wells3 as pure iron-carbon alloys has been measured the result shown in Fig. 15. Equally influenced rate of cooling and in are interested. this we structure, intensely ing

FUNDAMENTAL CHARACTERISTICS

appearance of solid-solution Fig. 16. Microscopic characteristic twinned grains. 1000X. (Vilella)

austenite.

21

Note

The decomposition of the solid-solution austenite does not begin instantly when its temperature is lowered to that at which, in time, it will transform. Instead, there is a definite period of lag that is pre sumably occupied by nucleus formation or the chance association of sufficient atoms of the new constituent to form a permanent crystal lite. At any rate, this reluctance is very definite and constant for any particular austenite, and a degree of undercooling is possible that, of course, depends on the rate of the cooling. Modes of Carbide Dispersion. When a new constituent develops within a metal, generally there is a wide range of final size of the in dividual precipitated particles, depending on the specific diffusivities involved; but in contour, the constituent generally conforms to one or more of the three primary categories — films, fllaments, or parti-

ALLOYING ELEMENTS IN STEEL

22

Fig. 17. Transformation to pearlite in its beginning. Note nodular growth from mation

of both ferrite

and carbide.

a eutectoid steel arrested at nuclei and simultaneous for 1000X. (Vilella)

cles. Depending on circumstances, iron carbide may form in plates or lamellae, intergranular cells, or somewhat spherical particles, over a range of sizes. Nowhere in metallurgy is there a more complete rep resentation of the effect of the distribution of hard particles in a soft plastic matrix than in steel. The finer the dispersion, the greater the hardness of the aggregate. Thus, the extraordinary range of properties is easily accounted for. In the clusters of lamellae, the films of hard carbide may be as thin as 0.0000005 in. (125 A) or as thick as 0.00008 in. (20,000 A). When particles are spheroidal, size range is from only a few Angstroms to nearly 0.001 -in. diam and the arrangement of the dispersed system is under control through the cooling and reheating of the steel — in short, heat treatment.

FUNDAMENTAL CHARACTERISTICS

to pearlite arrested when Fig. 18. Transformation tenite had transformed. 1000X. (Vilella)

25%

23

of the aus-

Consider for simplicity a steel to C. Heated to about 1400 F, the steel consists 0.70 containing 0.90% wholly of austenite in polyhedral grains, as in Fig. 16. On slow cool ing at a rate of 4 or 5° F per min, the austenite reverts to ferrite and The Lamellar Family of Structures.

carbide with the lamellar arrangement. The undercooling will amount to 15 to 20° F and the transformation begins at 1300 to 1305 F. The temperature of the steel may actually rise slightly during the trans formation, while the final reaction may even occur at a lower tem perature. The lamellae are spaced about 0.000015 in. (3750 A) apart and are about 0.000002 in. (500 A) thick. The hardness is about 210 Bhn. The mode of transformation may be seen in Fig. 17 to 21. The effect of increasing the rate of cooling a 0.75% C steel may be seen in the following table:

ALLOYING ELEMENTS IN STEEL

24

to pearlite arrested when 50% Fig. 19. Transformation tenite had transformed. 1000X. (Vilella)

1400

Cooling rate, to transformation 5° F per min 230° F per min 1800° F per min

F

In another instance, 1400

Cooling rate, to transformation

F

F per min F per min 400° F per min 3° 100°

a 1.0%

of the aus-

Transformation temperature, F

Brinell hardness number

1305 1240 1010

210 315 415

C

steel showed the

following behavior:

Transformation temperature, F

Brinell hardness number

1300 1260 1245

217 262 302

With still more rapid cooling, the fine lamellar structure ceases to form, and with even 1.1% C, the hardness of the lamellar structure

FUNDAMENTAL CHARACTERISTICS

to pearlite arrested when 75% Fig. 20. Transformation tenite had transformed. 1000X. (Vilella)

25

of the aus-

about 444 to 461 Bhn at this limiting condition. As Mehl4 so aptly points out, at the rate at which the lamellae are becoming thin ner with lower temperature of formation, a temperature must soon be reached at which each lamella would be of smaller dimension than the supposed crystal unit of cementite. Figures 22 to 26 are micro graphs of these lamellar structures for different temperatures of for mation, each having the hardness indicated. For air cooled sections to % in. thick, the hardness varies with carbon content about about as shown in the middle curve of Fig. 32. When the rate of cooling is too rapid for the lamellar type of trans is

formation at about 1000 F (550 C), the austenite does not transform until a very much lower temperature is reached, assuming the cooling medium itself to be at a temperature below about 200 F (100 C) ; the

26

ALLOYING ELEMENTS IN STEEL

Fig. 21. Transformations, shown in Fig. 17 to 20, when all austenite had just transformed to pearlite. 1000X. (Vilella)

transformation is then of a different sort and a much harder con stituent, martensite, is formed. The rate of cooling that just prevents the reaction in the vicinity of 1000 F and results in the lower tempera ture reaction is called the critical cooling rate or critical quenching rate. The former transformation and its temperature are sometimes spoken of as Ar', and the second as Ar" or more commonly now as Ms; the critical rate is the cooling rate separating the two. Many steels, however, cooled at about the critical rate, transform in part by one mode and in part by the other, even in the same grain, reflecting the fact that nucleation is properly expressed as a probability. Extremely slow cooling of many of the steels under discussion may not yield the well-formed lamellar structure at all, particularly if they are low in all alloying elements, the austenite is fine-grained, or some

FUNDAMENTAL CHARACTERISTICS

Fig. 22. Coarse pearlite formed at about ness: Rockwell C 5, 170 Bhn. 2500X. (Vilella)

1325

F

(720

27

C). Hard

is

is

is,

carbide is left undissolved in heating. Instead, the carbide is partly spheroidized or somewhat cellular, i.e., at grain boundaries. For this rea son, it is incorrect or misleading to show in a general diagram a trans formation at equilibrium temperature of austenite to lamellar pearlite, because pearlite was the name suggested by the pearly appearance of the lamellar association. The Proeutectoid Reaction. For compositions somewhat removed the composition which trans from the eutectoid composition (that to austenite at the lowest forms wholly possible heating temperature), a rejection of ferrite or carbide may occur before the austenite, thus restricted in carbon, transforms to the lamellar structure. The maxi limited by mum possible amount of such proeutectoid constituent controlled but the actual the composition, by the rate proportion

28

ALLOYING ELEMENTS IN STEEL

Fig. 23. Coarse pearlite formed at about 1300 F (705 Rockwell C 15, 208 Bhn. 2500X. (Vilella)

C).

Hardness:

is

is

is,

of cooling. The higher the rate of cooling, the smaller the amount of the proeutectoid ferrite or carbide to be found. Thus a 0.50% C steel is readily cooled so that a very small amount of ferrite separates be fore the remaining austenite transforms to a fine lamellar structure. Somewhat the same state of affairs exists for the hypereutectoid steels with respect to carbide, which however, more likely to separate than ferrite for an equal remoteness in composition from the eutecform in the grain toid. In noneutectoid steels, both constituents ferrite becomes somewhat more of the but boundaries austenite, irregular than the carbide envelopes. The proeutectoid reaction essentially one of diffusion. Sizable cannot form unless carbon can migrate there regions of ferrite alone function of temperature. Thus, the a from and such diffusivity

FUNDAMENTAL CHARACTERISTICS

24. Medium pearlite formed at about Rockwell C 30, 296 Bhn. 2500X. (Vilella)

Fig. ness:

1225

F

(665

29

C). Hard

the degree of undercooling, the smaller the size of ferrite or masses formed; considering that the entire time within the range at which any ferrite forms is often only a matter of minutes, it is not surprising that equilibrium with respect to the proeutectoid reaction is not reached in most instances. It is interesting that some proeutectoid ferrite forms in hypoeutectoid steel even when the austenite is brought unchanged to temperatures where austenite is wholly unstable. This is a good illustration of a principle that is exemplified frequently in the study of transformations. The reaction occurs that, for one reason or another, can take place most rapidly and it often has little to do with the most stable end product. Here ferrite can nucleate fairly easily, but once pearlite has begun to form, the rate of growth of pearlite is so much faster than that of ferrite that it greater

carbide

ALLOYING ELEMENTS IN STEEL

30

Fig. 25. Fine pearlite formed at about 1200 Rockwell C 33, 319 Bhn. 3000X. (Vilella)

F

(650

C).

Hardness:

"absorbs" the remaining austenite. An interesting corollary of that the composition of pearlite is rarely eutectoid in noneutectoid steel. The properties, however, do not depend on composi tion so much as on the carbide spacing which is a function of the temperature of formation of the pearlite. Acicular Mode of Transformation. When the eutectoid carbon steels are quenched at a rate exceeding the critical cooling rate in a medium near room temperature, the transformation is rapid and starts to occur at a temperature near 450 F (230 C); the product is hard and more voluminous than the other structures. This structure is called martensite and is illustrated in Fig. 27 to 29. Microscopically it has a characteristic acicular structure — evidence that its mode of formation is of successive nearly instantaneous transformations on rapidly this

is

FUNDAMENTAL CHARACTERISTICS

Fig. 26. Fine pearlite formed at about Rockwell C 40, 377 Bhn. 3000X. (Vilella)

HOOF

(595

C).

31

Hardness:

crystallographic planes of the austenite crystal. The transformation, not involving diffusion, occurs by the so-called "shear" mechanism; the orientations in adjacent plates from the same austenite grain are not identical because the transformation may proceed on different austenite

untransformed austenite must is therefore equivalent in metal. Its hardness can on this and is not yet com circumstance, however, scarcely depend understood. pletely Nature of Martensite. The atomic structure of martensite is shown schematically in Fig. 30. The carbon atoms are ordered preferentially in the octahedral sites at the center of an [001] edge or a (001) face. Since these tend to distort the lattice unidirectionally, an x-ray difplanes

and

the remaining

suffer some deformation. The martensite part to a fairly fine flat-grained strained

ALLOYING ELEMENTS IN STEEL

32

Fig. 27. Microscopic 2000X. (Vilella)

appearance

of martensite

fraction pattern will show a tetragonality the carbon content (Fig. 31).

in

eutectoid

steel.

that is a linear function of

is

a

is,

It is interesting to note that the carbon atoms in austenite, which are distributed at random among three crystallographic directions, are by the very nature of the shearing reaction that produces martens ite transformed without diffusion without changing their (that nearest neighbors) into an array that distorts the martensite crystal along single lattice direction, the c-axis. The reason for the hardness of martensite has been discussed for at least 50 years without any really satisfactory explanation. The lat tice certainly highly strained, but this does not seem sufficient to explain all of the hardening. Very recently, direct visual confirmation of the hypothesis that profuse multiple twinning takes place during

FUNDAMENTAL CHARACTERISTICS

33

Fig. 28. The appearance of martensite at low magnification in a fine-grained steel. Note that the acicular structure evident in Fig. 27 and Fig. 29 is much less obvious at this magnification (100X). The austenite grain size cannot be clearly distinguished. This is a slightly hypereutectoid steel and some undissolved carbides are evident.

martensite formation has been obtained by electron transmission photographs of very thin films. These twins are between 10 and 100 A thick. So far, however, it is not clear how, if at all, these twins could so profoundly influence mechanical properties, notably hard ness.

The relation of the carbon content to the full hardness of martens ite is shown in Fig. 32. Note that the high-carbon region of this curve is broad. Because of an increasing tendency for some of the softer austenite to be retained permanently through the quench, highcarbon steels often show a reduced hardness. When quenched, how ever, with some carbide undissolved, they reflect the hardness effect of residual hard particles as well as that of the more nearly pure mar tensite. The highest hardness figures in the range represent a proba ble maximum for quenched steels, in the approach to which the metal may be refrigerated to induce further transformation of any retained austenite.

ALLOYING ELEMENTS IN STEEL

34

Fig. 29. Martensite plates (slightly tempered during mounting in Bakelite) in a 1.7 wt% C steel, in a background of retained austenite resulting from the low Mf of this steel (100X). Some un tempered martensite is also evident; it may possibly have formed isothermally at room temperature (a rather rare phenomenon).

Tempering. When quenched martensite in carbon steels is reheated for 1 hr at various temperatures, the hardness follows the pattern shown in Fig. 33. This behavior is quite general for all carbon steels. Apart from a slight initial hardness increase in steels containing more than about 0.8% C, the hardness falls off continuously. In the initial stages of reheating — to perhaps 300 F (150C) for periods up to 1 hr — the martensite precipitates a hexagonal closeFe2.4C, leaving a matrix which is packed e-carbide, approximately very nearly cubic martensite but, since it may contain 0.25% C, is still supersaturated. If the carbon content of the steel is sufficiently high, the e-carbide may provide enough dispersion strengthening to compensate for the less hard low-carbon martensite. It is this that accounts for the initial increase in hardness at low tempering tem perature. At higher temperatures, beginning at about 480 F (250 C), the £-carbide and low-carbon martensite react to form Fe3C (as thin

II

I

!■

FUNDAMENTAL CHARACTERISTICS

0

35

Probable positions for carbon atoms

Fig. 30. Body-centered tetragonal structure of martensite showing possible (equivalent) sites for carbon atoms. Not all these possible sites for carbon are occupied in martensitic steel.

platelets) and cubic ferrite only slightly supersaturated with carbon. As the temperature is increased, the Fe3C particles grow larger and more nearly spheroidal, consequently contributing less and less to the strengthening of the material but creating a product whose ductility is very much improved over that of the as-quenched martensite. Spheroidal versus Lamellar Structures. There are two means for securing final hardness between 190 and 400 Bhn in higher-carbon steel. The hardness may be developed by cooling the appropriately heated and substantially austenitic steel at a rate that will produce a lamellar structure of the desired fineness and hardness, or the heated steel may be quenched to form hard martensite and then tem pered to develop the same hardness. While it might be supposed that there would be no great difference in steels of the same hardness re gardless of the heat treatment selected, actually the mechanical properties are quite different depending on whether the carbide is in the form of lamellae or spheroids. This is shown in Fig. 34, which compiles the results of a series of tests on specimens heat treated as

2.80

0.2

0

0.8

0.6

0.4

1.0

1.2

1.6

1.4

Carbon, % Fig. 31. Tetragonality

outlined

above.

of martensite as a function

This carbon

steel

contains

0.23% Mn.

of carbon content.

0.84% C, 0.20% Si, and

is

is

is

is

is

is

it

is

is

is

is a

is is,

For any hardness developed within the range, the plasticity values of the tensile test are superior for the spheroidized structures, that for the quenched and tempered steel. Only the tensile strength lower for the quenched and tempered steel — commonly called "the heat treated" steel. This superior strength in the pearlitic series reflection of the greater cross section at the beginning of local elonga tion and the termination of general elongation, for the reduction of area far greater for the tempered steel. The significance of greater reduction of area this: at the same hardness, the steel with the still capable of greater plastic flow spheroidal carbide dispersion even though has considerably greater notch effect at the zone of local constriction. This property highly valued in many applica tions. It should be noted that the numerical value of elongation because the length-to-diameter ratio not standard but low; this about 10:1. Figure 34 shows in a quantitative way the value of heat treatment to develop the carbide dispersion of tempered martensite or bainite. If a steel incapable of acquiring these structures in any large sec

FUNDAMENTAL CHARACTERISTICS

37

1100

1 11

1000

900 800

-

»700 w

65

60

dartensit ic struct jre (quer ched)

\

a. E

33^

f

-

A

\50%

600

O

33

£

38

f|

42 50

-Ms

55

Mso

400

M90

0.59 %C

0%

200

0.5

a

X

1

1

1

2

5

1

1

10

1 102

1

1 I Hr

1Min

1

1

l04

103

Ni

I

1

l05 1 Day

1 106 1Week

Time, sec

Fig. 113. Isothermal transformation rates for a carbon steel com parable with that of Fig. 112 except for nickel. Time intervals are on logarithmic scale.

fective proportionally as would be expected; indeed, they are no more In Fig. 117,20 this is illustrated for a effective than chromium. 0.35% C steel in which the molybdenum was substantially dissolved. The hardness distribution curves indicate that little pure fully hard martensite existed in the sections, but that the manifestations of Indeed, a fairly reliable hardenability are apparent notwithstanding. can be made by certain characteristics of estimate of hardenability the transformation at the relatively high temperatures that are remote from those producing martensite. Depression of Ar' a Reflection of Hardenability. When any par ticular type of steel is cooled at a specific rate that depresses the transformation to 100 to 300 F below the Ae! temperature, its precise undercooling is dependent on the same fundamentals that determine its hardenability. The undercooling then becomes a measure of hard enability, but a temperature should be chosen from which to estimate

EFFECTS OF ALLOYS IN HARDENING

// /

U— 0.5%

jl—

1

2

/ 6.5%Mnv/^

2.2% Mn^ ? go/ ii.

— 1/— /

/S x

#

Mn

1.0% Mn

5 10 20 Sec

'

f

980F

I/

5 10 20 Min

/

155

a.r/oMn

l3.4%Mn^*~-

5

1

Hr

10

1

2 5 Days

10

4 Weeks

Fig. 114. Individual transformation curves, from 100% austenite, at 980 and 595 F in a series of manganese steels with about 0.55% C. Time intervals are on logarithmic scale.

undercooling. For the

same nominal steels, this may be regarded as therefore ignored entirely, only the temperature at which cooling halts being observed. When dissimilar steels are com pared, the halt point on slow heating or the temperature inducing the first austenite may be employed as the reference for undercooling. This is illustrated in any of the three cooling curves of Fig. 118 (alike for all three steels), which are intersected at the transformation temperature, in terms of undercooling. The curves are for three differ ent chromium steels of increasing hardenability. The degree of un is seen to reflect the of the additions dercooling deep-hardening element. Also, the relative hardness, though not generally of martensitic range for these steels, reflects hardenability, because the products are formed at lower temperature and are thus harder. Perhaps such observations of the depression of the Ar7 temperature are reflections, rather than measures, of hardenability, as are also hardness deter minations of the products of slow cooling, but when employed for a

constant

and

156

ALLOYING ELEMENTS IN STEEL

single type of steel by an experienced investigator, they are useful in predicting hardenability. Dissolved Alloy Versus Hardenability. It would be helpful to have some method for estimating the increase in critical size with increase in alloying element and carbon in steels of uniform grain size and homogeneity of the austenite. The band in Fig. 119 shows the in fluence of manganese in increasing the hardenability (critical size) of 0.45% C steels. The quench is a vigorous movement in water with an approximate value of H = 5. The wide band is employed not so much because the data are inconsistent as because there are not enough data based on closely corresponding compositions and grain sizes. Several attempts have been made to calculate harden ability from composition. The efforts are based on information such as that in Fig. 119. For a fixed grain size, the effect of manganese on a measure of hardenability such as Di is dependent on composi tion. Attempts have then been made21 to use a formula such as: . fMn . fsi . fs . fp . fcr . . . etc, where the base Di D! = (base is a function of carbon content and grain size, and fx is the "multi plying factor" for an element X. Then fx is a function of composi tion. This equation assumes that each element behaves independently

EFFECTS OF ALLOYS IN HARDENING

157

l^j-in. rounds of four of hardness across Fig. 116. Distribution 0.35% C steels with chromium as indicated. The bars as quenched in oil were free from undissolved carbide.20

of hardness across l$4-in. rounds of three Fig. 117. Distribution 0.35% C steels with molybdenum as indicated. The bars as quenched in oil were free from undissolved carbide.20

of all others; to the extent that this assumption is reasonable, the equation predicts behavior moderately well. For only modest addi tions of many elements, however, agreement between the equation and experience is not good; a method such as this has value only as a last resort when experiments are not possible. At present, the picture is muddied further by disagreement between various workers as to the correct values of the multiplying factors. When present in relatively large proportion, the alloying elements

EFFECTS OF ALLOYS IN HARDENING

Undercooling

time,

159

min

Fig. 118. Effect of the same cooling rates on the transformation temperature of three steels of like carbon content; chromium contents of steels (2) and (3) were 1% and 3.2%, respectively. Degree of un dercooling rather than actual temperature is shown in the ordinate ; time in the range below critical temperature is reported on the linear scale.

contribute relatively less to hardenability if used singly. The most ef fective additions are combinations of elements, as would be suggested by the equation on page 156 — for example, the nickel-chromium or steels. nickel-chromium-molybdenum Elements in Carbides. It has been shown that the proportion of an element that becomes dissolved in austenite is that which is effective in increasing hardenability. The portion of an alloying element com bined in undissolved carbide, if not inert, surely plays no helpful part in hardenability, but it has already exerted an influence in the opposite direction because it has restricted grain growth and because it may act as a powerful nucleation site for pearlite. It is well to re peat that the heating temperature and interval exert a very strong influence upon this partition of the carbide-forming elements between austenite and the inert undissolved carbide. The heating schedule determines which role is dominant for these powerful dual-natured elements. Vanadium is the best element to use in exemplifying the carbide effect in hardening, for its influence is more intense than, although similar to, those of tungsten and molybdenum. Titanium is a stronger

ALLOYING ELEMENTS IN STEEL

160

0.4

0.6

0.8

1.2

1.0

Manganese,

1.4

%

Fig. 119. Influence of dissolved manganese content on critical size of rounds of 0.45% G steels. For water quench, H = about 5. Slope of band is probably representative.

carbide-former than vanadium, but with appropriate compositions and temperatures, its action is comparable in principle. The undissolved carbide of vanadium steels robs the austenite of some of its carbon content and restrains grain growth so that, at low heating temperature, it generally lowers hardenability. Figure 120 shows the grain-growth characteristics of two 0.90% C steels, C and V, which are substantially identical except for 0.27% V in steel V. (The extraordinary fineness of the austenite grain requires an exten sion of the ASTM grain size numbers beyond No. 8; the same sys tem of numbering is maintained, There is a zone in however.) which both steels coarsen at maximum rate with increasing heating temperature.

Consider now the hardness distribution across diameters of the two

steels as quenched from successively higher temperatures, as shown in Fig. 121. 22 When the plain carbon steel has achieved a grain size of No. 4 to 5, its hardenability is almost precisely equal to that of the

5,

4

8

7

is

F

is,

vanadium-bearing steel when its grain size is No. 7 to 8. At the tem 1650 perature inducing this grain size in the vanadium steel, that the C), vanadium-rich carbide to be dissolved. just beginning (900 Accordingly, the first significant solution of the vanadium in the aus tenite has a deep-hardening effect equivalent to a change in grain size from or to or three ASTM numbers. At the heating tem

EFFECTS OF ALLOYS IN HARDENING

161

10

il

c

V

5

km 0)

3

D

^Steel

C

O

S

1400

2000

1600 1800 Heating temperature, F

Fig. 120. Approximate trends of mean grain size in the austenite of two 0.90% C steels alike except for 0.27% V in steel V. Note coarsen ing ranges. (From data of Zimmerman, Aborn and Bain)

at which a further increment of vanadium is dissolved, that to and (975 C), the steel has acquired a grain size of has become exceedingly deep-hardening and, in this respect, exceeds the plain carbon steel coarsened to a grain size of to Actually, a further solution of vanadium (more readily brought about in somewhat lower-carbon steel) induces an extraordinarily high hardenability when the grain size approaches or exceeds that of carbon steel indeed, bars larger than in. must be employed to permit estima tion of the increment in hardenability. Care must be taken not to add too much vanadium or columbium, or may be impossible to produce austenite. 100% 5

it

1

a

2.

1

4

F

1800

;

is,

perature

162

ALLOYING ELEMENTS IN STEEL

It is possible that vanadium may restrict grain growth by virtue of other mechanisms than its contribution of a particular kind of per sistent fine-carbide dispersion. It would be difficult to prove that there was not a dispersion of particles of an oxide or nitride that evades detection. However, the close parallel between marked grain growth and carbide disappearance in these steels as in tungsten and molybdenum steels (and, indeed, carbon steels) makes the present hypothesis satisfactory if not conclusive. Considering the preservation of fine grain after high heating temperature and the accompanying toughness in heat treated vanadium steels, it is perhaps proper to suggest the use of higher heating temperatures for vanadium steels and similar steels than are often applied. The advantages will be seen also in the subsequent observations on tempering. To sum up, the carbide-forming elements are natural grain-growth restrainers so long as some proportion of the fine carbide particles ele remains. As dissolved, they are very powerful deep-hardening ments. Advantage may be taken of both characteristics by a suitable heating schedule. Anomalous Hardening at Low Heating Temperatures. To employ a valuable carbide-forming element only as a source of hard inert par ticles is justifiable when abrasion resistance is needed as in cutting tools, provided sufficient carbon and alloy are present to form an

EFFECTS OF ALLOYS IN HARDENING Tungsten

163

steel

90

55 1

80

50 70

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45

60

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25

>

10

40

20

60

80

20 100 Hours at 1380

40

60

80

100

F

0

20

20

0

i

a

Fig. 122. (Left) Change in hardness developed on quenching 0.55% C, 2% W steel, as a function of heating time at relatively low temperature. (Right) Change in proportion of austenite.23

a

is

C

A

F

is

a

a

is

a

austenite of high hardenability and furnish needed strength in the resulting tempered martensite. However, when heating temperatures are employed that are so low as to develop only low-carbon lowalloy austenite composition, the final properties are rarely such as would be demanded for any practical purpose. In so alloying steels with tungsten, molybdenum, vanadium, and to some extent chro mium, poor use made of the alloy and some rather surprising be haviors may result. Consider 2% W steel with about 0.55% C. At equilibrium, such composition actually composed of ferrite, tungsten-rich carbide and a little austenite at a temperature of 1380 (750 C). This has not always been known. In the course of developing information on equilibrium conditions, such a steel was heated for a number of in short heating followed by creasing intervals at this temperature.23 52, but increasing the quenching produced a hardness of Rockwell heating time did not increase the resulting hardness; instead, longer heating gradually reduced the as-quenched hardness. The behavior shown in Fig. 122. Having regarded the increase of heating period as a means of dissolving more carbon and alloy, this was perplexing

ALLOYING ELEMENTS IN STEEL

164

it

is

is

it

is

a

is,

result. That it was not a matter of decarburization was verified, and in the course of microscopic observation it was found that the pro portion of austenite formed at the heating temperature was actually falling off with lapse of time. The unexpected behavior demanded an explanation and it became quite apparent that the unique factor was the low heating temperature. Presumably, at this low heating temperature, carbon migrates with moderate velocity in terms of microscopic distances. Accordingly, at first, a considerable proportion of austenite of rather low carbon content is formed. Later, much later, the tungsten migrates and asserts its influence to form ferrite, with a little austenite of higher carbon content and some tungstenrich carbide. The carbide-forming elements, when actually dissolved, exert a far greater effect on the formation of pearlite than do corresponding proportions of the manganese, nickel and silicon. Additions of the carbide-forming elements do not preserve much austenite after the quench. Chromium in solution fosters some austenite retention but far less than corresponding amounts of manganese. While the effect of the undissolved carbide-forming elements on grain growth is marked and the hard residual carbides have special value in tools, the most interesting of their functions is revealed only in tempering when they are initially in solution in martensite and are precipitated in a rather special manner. Elements in Nonmetallic Inclusions. The influence of nonmetallic inclusions is presumably much the same in alloy steels as it is in ironhow carbon alloys and this effect has already been discussed. It ever, obscured in the presence of carbide-forming elements. In the discussion of carbon steels, particular attention was given to the very effective additions of aluminum that act through dispersion of alu minum nitride inclusions. The aluminum treatment, when skillfully utilized, applicable to the elevation of grain-coarsening temperature in practically all steels other than high-aluminum or rimming com positions and has been so used in many compositions. worth repeating that the dispersion, whatever It may be, of an aluminum compound may act not only to restrict austenitic grain growth and thereby lessen hardenability at normal heating tempera tures, but also may act directly by stimulating nucleation. This evident in typical "fine-grained" steels (with high coarsening tem perature) that have had their austenite grain greatly coarsened by excessive heating. Under certain circumstances, such coarsened steels have the low hardenability of fine-grained steels with their tough centers. This has been described by Bain and need not be considered in detail here. However, may not be out of place to consider the magnitude of the change in hardenability brought about by coarsen ing of the austenite grain size under circumstances in which no other influences are significant — such as homogeneity of austenite, undis solved carbide, particle nucleation. Figure 123 shows the effect of

EFFECTS OF ALLOYS IN HARDENING austenitic

critical

165

grain size on the absolute hardenability in terms of ideal for four steels of the following approximate analyses:

size

1.

Mn Mn C, 0.30% Mn C, 0.30% Mn

0.40% C, 0.80% C, 0.90%

2. 0.42% 3. 0.75% 4. 0.90%

The falling off of hardenability of steel 3 at finer grain size is thought to be abnormally great because of undissolved carbide resulting from inadequate heating. While these examples are not as uniform as might have been hoped for, they indicate in a general way that in such steels a decrease of three in grain size number may alter the ideal critical steels size by about 0.25 in. Nickel, chromium, or higher-manganese been availa have been used had suitable data as illustrations might ble, but in steels wherein carbide solution and subsequent homogenization are difficult, the single factor of grain size would not have been free from other influences. Transformation rates at higher tem perature for different austenite grain sizes are shown in Fig. 124. The nonmetallic inclusions that are beneficial in steel are those that aid in retaining a fine austenite grain size or nucleate martensite formation and thereby contribute toughness. It is of no great benefit if they merely reduce hardenability. The characteristics of such nonmetallics are these: 1.

2. 3.

of easily eliminated inclusions in the liquid steel, probably by acquiring large size. Relatively high solubility just below the solidification tempera ture and low solubility at heat treating temperature. At the temperature of falling solubility, low diffusivity of at least one of the elements entering into the compound of the Formation

inclusion.

4.

A capacity for stimulating austenite nuclei on heating or possi bly martensite nuclei on quenching.

These characteristics probably insure the absence of injuriously large inclusions and provide the properties that are regarded as valu able and to which the name "body" was at one time applied. Some agents appear to act in these ways (in addition to the carbide par ticles that are in part similar), although there is question as to whether they are quite as effective as aluminum. Microscopic Homogeneity and Hardenability. As has been shown, the solution of carbide is a time-consuming operation; even though it is rapid in pure iron-carbon alloys, it is slow in some alloy steels. This is because diffusion itself is time-consuming, and only through diffusion may austenite be rendered homogeneous, for the austenite formed at the former location of carbide particles is very rich in car bon and alloy. Frequently, in quenched steels that have been heated for hardening from a coarse pearlitic state, the markings of the origi

166

ALLOYING ELEMENTS IN STEEL

12

4 6 5 ASTM austenite grain size 3

Fig. 123. Ideal critical size in four compositions tenitic grain size at quenching temperature.

as changed

7

8

by aus-

nal lamellae are clearly visible in the microscopic specimen of result ing martensite. The reason the etching agent reveals the pattern is that the martensite is of uneven carbon content: low in the former ferrite lamellae zones, high in the cementite lamellae locations. Lowcarbon martensite, because of its high Ms temperature, is partially tempered even during a quench24 and consequently etches somewhat more rapidly, that is darkly, than a higher-carbon material. An inhomogeneous austenite is not unlike an intimate mixture of two steels (and intermediate ones also). The high-carbon and alloy regions of microscopic extent have high hardenability, the remainder

Fig. 125. Hardness distribution in 1-in. rounds of eutectoid steel as influenced by structure entering heating bath. Constant heating sched ule. (Left) Short time at 1400 F; (right) longer time at 1500 F.25

ALLOYING ELEMENTS IN STEEL

168

126. Hardness distribution in quenched round bars of SAE steel. Vigorous water quench from 1525 F. (Courtesy W. M. Lindsey and E. L. Roff, South Chicago Works, United States Steel

Fig.

1040

Corp.) Fig. 1040

127. Hardness distribution in quenched round bars of steel. Average oil quench from 1525 F. (Carilloy Steels)

SAE

low. For pure martensite, the steel must be quenched at the required critical rate of the leanest austenite. For practical purposes, hardenability is impaired by inhomogeneity resulting from insufficient heat ing; this is a good example of the "weakest link in a chain." Clearly, the simplest way in which to render homogenization easy is to form austenite from the finest possible dispersion of carbide com patible with other requirements, such as machinability. Thus, the structure prior to heat treatment, especially when minimum heating time and temperature are employed, has a marked influence on hardenability. Double heat treatment, from high temperature for homoge nization and then at lower temperature for fine grain size and toughness, is exceedingly desirable and recommended practice. The influence of the structure entering the heating bath (for a constant schedule) on hardenability is illustrated in Fig. 125. A short heating period (left) fails to equalize the three conditions of the steel, whereas a longer interval (right) at high temperature nearly does so. The three structures were a normally coarse spheroidized state, a more finely spheroidized one and a fine pearlite produced by a mild oil

EFFECTS OF ALLOYS IN HARDENING 3-in. t-m.

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portion of retained austenite and this view certainly compatible with the results. At any rate, the scatter for the several steels greatly reduced when the steels are so tempered as to eliminate austenite present after the quench. Loss of Toughness in Intermediate Tempering

Low-Temperature

toughness of hardened

plasticity and notch Increasing by tempering has thus far been spoken

Tempering. steel

200

250

300

350

400

Hardness,

450 Brinell

500

550

600

Fig. 184. Range of tensile properties in several quenched and tem pered steels at the same hardness values. (Janitzky and Baeyertz)

of

as a continuous effect, becoming more marked as the tempering temperature is raised. This is not strictly true, for carbon steel and many low-alloy steels do not always acquire augmented notch tough ness after tempering at a specific low tempering temperature. The elongation in the tensile test generally increases slowly with any tem pering increment, but in the vicinity of 500 to 600 F the tempering treatment often results in a notch toughness lower than that of steel

200

250

300

350

400

Hardness,

450

Brinel

500

550

600

1

Fig. 185. Range of values of reduction of area in several quenched and tempered steels at the same hardness values. Note spread at high (Janitzky and Baeyertz) hardness.

tempered at 300

F or scarcely greater than that of untempered

steel.

Higher tempering temperatures usually raise the notch toughness to much higher values. The curves of Fig. 186 are typical of many steels.

EFFECTS OF ALLOYS IN TEMPERING 200

300

400

Temperature, 300

227

C

400

500

600

Tempering

temperature,

Fig. 186. Loss of notch toughness in Izod after tempering at about 600 F.

700

800

900

F

test on several

alloy steels

The phenomenon of an unexpected lowering of notch toughness, or the inability to undergo deformation under adverse conditions, has

the subject of much discussion. In 1924, Grossmann9 offered an of this reduced toughness in steels tempered early explanation at about 550 to 600 F. Some of his data are presented in Fig. 187. The upper chart shows the characteristic minimum notch toughness, and below is shown the length change accompanying tempering. An anomalous retardation of shrinkage corresponds to the lowered toughness. This failure to contract uniformly with increased temper been

ALLOYING ELEMENTS IN STEEL

228

0.0005

-0.0010

o§

0.0015

0.0020 1 Hr

0.0025

0

100

200

300

400

500

600

Tempering temperature,

700

800

900

F

Fig. 187. Relative notch impact values (top) and change in length (bottom) resulting from tempering at various temperatures. Note ef fects at about 600 F. (Grossmann)

ing amounts to a definite expansion in the same steel when the heat ing for hardening is at a higher temperature such as 1800 F. These circumstances point to the preservation of austenite on quenching and its subsequent transformation during tempering. The same temperature range of tempering has been found to af fect adversely the capacity of the steel to flow in shear during a tor sion test. Luerssen and Greene10 observed a much reduced angle of

EFFECTS OF ALLOYS IN TEMPERING 0

100

Temperature, 200

Drawing

229

C

300

400

temperature, F

Fig. 188. Irregularities in property changes accompanying at about 500 F. (Luerssen and Greene)

tempering

twist in their torsion impact test (Fig. 188). Cohen and co-workers11 have noted that the first Fe3C to form during tempering, which is formed in this temperature range, is essentially two-dimensional; they suggest this morphology may be responsible for the various effects around 575 F. Recently, in vacuum-melted steels made from the purest available starting materials, Gensamer12 and Capus13 have shown that the loss of toughness is absent and suggest that some impurity is respon sible. Capus believes that antimony, arsenic, tin, or phosphorus may cause the embrittlement. At the moment, it seems possible that the effect may be due to more than a single cause. A further interesting point in this connection is the nearly complete freedom of austempered specimens from this kind of brittleness in the hardness range that manifests it in quenched and tempered specimens of the same steels. This is illustrated in the work of Payson and Hodapp14 (Fig. 189). Heat Evolutions in Tempering. One useful method of experimen tation confirms the views on tempering previously expounded. The

ALLOYING ELEMENTS IN STEEL

230 60

SAE

3240 from 1500 F

/" 000

(Numt >ers refer t