United States Department of Commerce

Technology Administration

National Institute of Standards and Technology

NISTIR 3992

STRUCTURE-PROPERTY RELATIONSHIPS IN

MICROALLOYED FERRITE-PEARLITE STEELS

PHASE 1 : LITERATURE REVIEW, RESEARCH

PLAN, AND INITIAL RESULTS

P.T. Purtscher

Yi-Wen Cheng

—QC—

100

.U56

#3992

1993

NISTIR 3992

STRUCTURE-PROPERTY RELATIONSHIPS IN

MICROALLOYED FERRITE-PEARLITE STEELS

PHASE 1: LITERATURE REVIEW, RESEARCH

PLAN, AND INITIAL RESULTS

P.T. Purtscher

Yi-Wen Cheng

Materials Reliability Division

Materials Science and Engineering Laboratory

National Institute of Standards and Technology

Boulder, Colorado 80303-3328

July 1993

U.S. DEPARTMENT OF COMMERCE, Ronald H. Brown, Secretary

NATIONAL INSTITUTE OF STANDARDS AND TECHNOLOGY, Raymond G. Kammer, Acting Director
'
CONTENTS

ABSTRACT iii

1 . INTRODUCTION 1

1.1 Economic Factors 1

1.2 Application History 3

1.3 Outlook for Future Applications 3

1 . 4 Current State of Understanding

Strength-Toughness Relationships 4

1.4.1 Strength 4

1.4. 1.1 Effect of Grain Size 7

1.4. 1.2 Effect of Precipitation 7

1.4. 1.3 Effect of Phase Transformation 10

1.4. 1.4 Effect of Solid Solution Strengthening 13

1.4. 1.5 Effect of Texture and Dislocations 13

1.4. 1.6 Summary of Strength Effects 13

1.4.2 Hardness 15

1.4.3 Toughness 16

1.4. 3.1 Impact Testing 16

1 . 4 . 3 . 2 Fracture Toughness Testing 19

1.4. 3. 3 Tensile Testing 21

1.4. 3. 4 Fatigue Testing 26

1.5 Summary 28

2. PROPOSED RESEARCH PROGRAM AT NIST 30

2.1 Generation of Representative Microstructures 31

2.2 Characterization of Structure 35

2.3 Determination of Properties 35

2.4 Integration of Structure and Properties

for Modeling the Application 36

in
3. PRELIMINARY STUDY 36

3.1 Introduction 36

3.2 Material and Experimental Procedures 36

3.3 Results 40

3.3.1 Microstructure 40

3.3.2 Mechanical Properties 40

3.3.3 Fractography 44

3.4 Discussion 46

3.5 Summary 51

4.0 ACKNOWLEDGEMENTS 52

5.0 REFERENCES 52

6.0 APPENDIX A — Stress Analysis 58
ABSTRACT

The literature related to the relationship between micro-structure

and properties of steels after thermomechanical processing is

reviewed and a plan for research is outlined. The goal is to

provide a methodology which can predict the range of forging

parameters that will produce acceptable properties in the as-forged

condition. The first phase of the work will emphasize the microalloyed ferrite-pearl ite steels used to replace quenched and

tempered steels in forged automotive applications such as crankshafts and connecting rods. Emphasis is placed on predicting those

properties which are related to service life as a function of

process variables. In preliminary experiments on air-cooled, Nbtreated SAE 1141 steel, the tensile properties and ductile fracture

behavior were not strong functions of the hot-deformation

conditions. The conditon with the finest ferrite grain size

produced better impact properties. The tensile properties of the

steel can be estimated from the microstructure, but the difference

between predicted and measured values can be as great as 100 MPa.

v
.
1 . INTRODUCTION

Additions of small amounts of microalloying (MA) elements to a

basic C steel can produce properties in the air-cooled condition

which rival those of more expensive steels which require extra

processing (1-3) . The most common MA elements are V, Nb, and Ti,

but significant additions of Mn, S, Mo, and sometimes Si are also

used along with the common MA elements to optimize the beneficial

effect of MA. The following sections will cover background

information on MA ferrite-pearlite (a-P) steels.

1.1 Economic Factors

Ferrite-pearlite (a-P) steels, like the free-machining, medium C

grade SAE 1141, that have intentional additions of MA elements like

Nb, V, and/or Ti can be used to replace quenched and tempered (Q&T)

steels in applications where the steel is not subjected to severe

impact loading. The replacement can result in a cost savings of

approximately 30% (2). The cost advantages come in four areas: I.

lower alloy content, II. elimination of heat treatment cycle,

III. lower distortion, therefore no straightening or stress relief,

and IV. better machinability

.

Table 1 shows the composition for SAE 4140, a typical Q&T steel,

and SAE 1141 steel, a typical MA a-P steel. Q&T steels typically

contain 1.5 to 3% alloy content (Cr, Ni, Mn, and Mo) in the

composition so that the steel will have enough hardenability to

form martensite throughout the section during heat treatment. MA

a-P steels typically contain a slightly higher Mn content (1.0 to

1.5%) than a Q&T steel (0.7 to 1.0%), but no intentional additions

of Cr, Ni, and Mo. The other major difference in composition

between 4140 and 1141 is the S content. With 0.080 to 0.100% S

added to the MA steel, the machinability is enhanced compared to

that of a typical Q&T steel containing 0.025% max. S content (1,3)

.

Table 1. Chemical composition of typical Q&T and a-P steels

(mass %)

.

steel C Mn : S : P Cr Ni Mo Si MA

1141 0.42 1.49 0.099 0.015 0.06 0.06 0.01 0.26 0.04 Nb

4140 0.40 0.84 0.023 0.012 0.83 0.09 0.22 0.26 0.11 V

Figure 1 illustrates the different thermal cycles which are

required for MA and Q&T steels. The properties of a MA steel like

SAE 1141V are obtained by cooling the steel in air after processing. For a Q&T steel, the desired properties are only obtained

after an extra heat treating cycle which follows processing. The

1
Fig. 1. Schematic diagram showing the advantages in manufacturing

a crankshaft from MA a-P steels rather than Q&T steels (1)

relatively slow cooling that the MA steel receives produces less

distortion than the quench required for Q&T steels, and as a

result, there are no additional costs due to straightening or

stress relieving.

A recent study (4) has looked at a cost model for crankshaft

processing. The paper examined the replacement of Q&T alloy and C

steels (currently being used in production) with MA steel. The

model considered fatigue and machinability data for the various

steels. The MA steels offered a 10 to 20% reduction in cost, while

at the same time providing equivalent fatigue and superior machinability characteristics.

2
1.2 Application History

MA forging steels were first used in Europe in the mid-1970's (5-

7) . With controlled cooling, the Europeans found that they could

achieve a strength in a-P steels with V added that was equivalent

to tyical Q&T steels. The reductions in processing and alloying

costs were the initial incentive for replacing Q&T steels with MA

a-P steels, but later it became apparent that the machinability of

precipitation strengthened a-P steels was also superior to that of

Q&T steels, which provided further cost advantages to the MA a-P

steels.

The application of these MA forging steels has been aggressively

pursued since the late 1970's (7) . The toughness of the MA forging

steels was obviously lower than that of Q&T steels that had been

historically used for forging. Therefore, comprehensive component

testing programs were conducted by individual companies to form the

basis for the substitution of MA steels for Q&T steels. The

cumulative experience from component testing was positive; MA a-P

steels had adequate toughness for many applications like the

connecting rod in a car engine and car steering knuckle. The lower

toughness of the MA steels made it unacceptable for the front axle

beam in a truck. Subsequent studies in the U.S.A. (8,9) have come

to the same conclusion? MA a-P steels have adequate toughness for

many applications where Q&T steels have always been used.

A similar application for MA steels has been in the forging of

track links for crawler-type tractors (10,11). In this case, the

forging blanks are induction heated, forged, quenched directly from

forging to room temperature, and then tempered to the desired

hardness. The final microstructure is tempered martensite, but

fewer processing steps are required compared to the traditional

process where the steel is allowed to cool to room temperature

after forging, reheated, and then Q&T. Here, the MA additions are

for control of the grain size rather than for precipitation

strengthening

.

1.3 Outlook for Future Applications

MA a-P steel can be used in any application where moderate

strength, good fatigue resistance, and superior machinability are

needed. Direct replacement of Q&T grades with MA a-P steels

dictates that the steel be processed with existing equipment and

procedures while maintaining the strength of Q&T steels with an

acceptable toughness. A significant problem is defining what is

acceptable toughness. Q&T steels have been used in many

applications where toughness did not limit the design. Typically,

Charpy V-notch (CVN) impact testing has been used as the only

measure of toughness, and the lower impact toughness of MA a-P

steel compared to Q&T steel has made wider application of the

former difficult.

3
To improve the toughness properties of forging steels while

maintaining similar strength levels, bainitic alloys have been

developed with lower C content and higher Mn plus Mo for hardenability (12) . The nominal designations are AISI 1522 and 1535.

Higher MA additions are used to make up the strength that is lost

by lowering the C content. The lower C content does limit the use

of bainitic steels to applications that do not require surface

induction hardening, like automotive underbody parts (13) . Bainite, compared to a-P, has a finer structure and is more uniform

over the scale of a prior austenite grain.

Grassl et al. (14) have evaluated the properties of MA bainitic

steels as compared to MA a-P and Q&T steels. Initial results, Fig.

2, emphasize that Q&T steels have higher yield strength (YS) and

absorbed energy in room temperature CVN testing than MA bainitic

and a-P steels. The temperature transition curves for CVN absorbed

energy from the different steels are shown in Fig. 3 (15) . Again,

the Q&T steel (4140) appears much superior compared to the bainitic

(0.24 and 0 . 35C-Mn-Mo-V) and a-P (1045V) steels. The bainitic

steels are comparable to Q&T steel only at relatively high temperatures (>100 °C) . Significant changes to the MA steels will be

required before the air-cooled impact properties will match those

of Q&T grades.

1.4 Current State of Understanding Strength-Toughness

Relationships

1.4.1 Strength

From a physical standpoint, strength is related to the difficulty

of dislocation generation and motion. The harder it is to form and

move dislocations, the stronger the material is. Experimentally,

strength is evaluated by tensile testing and hardness measurements.

A tensile test provides precise information about yielding, strain

hardening, and fracture of a material, but is a destructive test.

Hardness is a more general parameter that approximates the yield

and ultimate strength of a material, and can be applied as a nondestructive test for inspection.

In the literature, equations to predict the mechanical properties

of a-P steel are numerous, but most are strictly empirical in

nature and are applicable to a relatively narrow range of compositions and processes. Modeling of strength by a single mechanism is

relatively advanced to the point where yield strength (YS) can be

described by the relevant physical process. Where multiple

strengthening mechanisms are operative, there is no generally

accepted model. Recent developments (16,17) have pushed modeling

towards a description of the stress-strain diagram.

Gladman and Pickering (18) and Pickering (2) have summarized the

literature regarding our understanding and prediction of strength

in a-P steels. A graphical representation of the effect of the

4
ROOM TEMPERATURE CVN (FT-LBS)

0 25 50 75 100

m

r

o

IA—

t

33m

2

Cl—

1

X

(A

Fig. 2. Plot of YS vs. CVN absorbed energy at room temperature for

the two types of steels (14)

.

TESTING TEMPERATURE (°C)

Fig. 3. Plot of CVN absorbed energy vs. temperature for four

different types of steels (15)

.

5
volume fraction of P on YS is shown in Fig. 4. Ferrite (a) grain

size and the volume fraction of pearlite f(P) are the most

significant factors influencing YS in a-P with no MA. YS varies

nonlinearly with P content (note that P has practically no effect

at fractions less than 0.2). To predict yield strength (YS) , the

general form of the eguation would be as follows

YS = ai + k

y

d“1/2

, (1)

where k

y

and a. are the constants from the Hall-Petch equation that

describe the effect of grain size, d, on the strength. The value

of d is generally taken as the macrograin size, but could be a

subgrain size under the right processing conditions. The a i term

is a catch-all for any strengthening mechanism that is operative;

the most frequent assumption used is that the effects of different

strengthening mechanisms are additive so that a. includes ass , a t

, ad' a

tx'

anci a

tr' where a

§s = solid solution strengthening, a t = precipitation strengthening, od = 'forest' dislocation strengthening, a

tx = a crystallographic texture parameter, and atr is the

strengthening due to phase transformation. The second part of Eq.

(1) , k *d’ 1/2 represents the grain size contribution to YS (a )

. y 9s

Pearlite Fraction

Fig. 4. Schematic diagram showing the contribution of different

strengthening mechanisms to the YS of steel as a function of

pearlite fraction (18)

6
1.4. 1.1 Effect of Grain Size

Hall-Petch relationships are important in understanding the change

in properties, but it can be difficult to employ Eq. (1)

quantitatively to a particular system. The measurement of grain

size can lead to problems. Even in an equiaxed and uniform grain

structure, there is some confusion regarding what is the effective

grain size. In most cases, the mean linear intercept, d*, of grain

size measured from the polished cross section of the material is

used. However, values of 1.75 to 2 times d* are sometines used to

account for sectioning effects. Frequently the definition of

'grain size' and sometimes its units are not even reported. These

factors contribute to the variation in the reported values of k

y found in the literature.

The problem of defining the effective grain size to use in a

predictive equation, like Eq. (1) , is acute when the structure is

not uniform or equiaxed. Low C steels were used in the derivation

of the Hall-Petch equation and d referred to the ferrite grain size

in the absence of a second phase. For higher C steels where there

is a large fraction of pearlite and the ferrite is present as a

network around the pearlite, the role of d in determining the

strength is more complicated. The thickness of the ferrite network

is the feature that is normally measured to include in Eq. (1) . Bainitic and martensitic structures in steels are inherently

acicular. Some processing conditions produce a duplex, but

equiaxed grain structure while others distort the grain structure

in the rolling direction (18) . It should also be noted that the function d' 1/2 seems to fit most

theories and a large body of data; however, equally significant

relationships can sometimes be obtained by using a factor of d to

the -1/4, -1/3, or -1 power (18).

1.4. 1.2 Effect of Precipitation

The effect of precipitation can be relatively large in a-P forging

steels; approximately 150 MPa increase in YS is not uncommon. The

precipitation occurs by MA elements combining with C and N, and the

maximum a t contribution could be estimated from the chemical

composition. In theory, a t can be modeled by considering the

influence of the particles on the movement of dislocations,

commonly referred to as dispersion hardening (19-20) . The

following equation can be used to describe the increase in YS

related to precipitation of V(C,N) particles:

A YS = (PPt) ]

1/2 l n

, ( 2 )

7
where f(ppt) = volume fraction of precipitates and x

ppt

is the mean

planar intercept of the spacing between particles in nm. K is a

constant which is nearly 6 N/m and b is the Burger's vector for

dislocations (approximately 0.25 nm)

Actual strengthening due to precipitation depends critically upon

processing. For example, precipitates containing Nb can form in

austenite during hot deformation, which retards recrystallization,

or the precipitates can form in ferrite during and/or after

transformation, which provides strength at room temperature. If

the precipitates form in austenite, they influence the final grain

size d, and therefore YS and CVN results, but not o . Data on the

variation of precipitation due to processing variables are found in

recent work by Leap et al. (21) . Precipitation is generally assumed to have nearly equal effect on

strength of a and P. Sage (5) found that this assumption was not

always true. In a steel with 0.35% C, 1.1% Mn, and a normal N

content of 100 ppm, precipitation was only found in proeutectoid a

and not in the pearlite phase. For other alloys, precipitates were

present in both phases.

The strengthening due to precipitation is usually inferred from

direct comparison with the predicted strength for a steel without

MA additions (5,22). Figure 5 shows the effect of V additions on

the ultimate tensile strength (UTS) of 0.45% C and 0.8% Mn steel

(5)

The variation in MA had no effect on the volume fraction of

ferrite f(a) and interlamellar spacing S

p

, so it was easy to

separate out the effects of MA and phase transformation. The a

increment increases with the V content at a rate of 100 MPa/ 0.1%

V. The scale of the precipitation strengthening (average particle

spacing L) , 0.03 to 0.1 /xm, is much finer than that which determines the YS of a and P.

The assumption that a adds to other mechanisms has not been

proven by fundamental studies in all cases. Theoretical models for

a t are based on the interaction between moving dislocations and

the precipitates. If the precipitates are coherent or semicoherent, the dislocations cut through the precipitates and affect

YS much more than the strain hardening and UTS. If the particles

are incoherent with the matrix, dislocations will normally bow out

around the obstacles and increase strain hardening and UTS as well

as the YS. For steels with <0.2% C, Gladman et al. (23) compared

the fundamental mechanisms and found the contributions of a t from

MA precipitates and of grain refinement, cr gs , were additive.

Hornbogen and Staniek (24) studied a due to Cu precipitation in

iron and how precipitation affected the sum of a t to a

gs

. In this

case, o t did add to cr gs

as long as the precipitates were coherent

(<10 nm m diameter) with the matrix. When the precipitates grew

larger and lost coherency, bowing of dislocations around the

8
VANADIUM, mass%

Fig. 5. Schematic diagram showing the contribution of different

strengthening mechanisms to the UTS of steel as a function of V

content (5)

.

9
precipitates was the dominant physical process affecting YS, and

the contribution of a was lost.

gs

1.4. 1.3 Effect of Phase Transformation

For the air-cooled, medium carbon a-P steels, the a tr contribution

is a result of pearlite formation. Transformation is typically

represented by a continuous cooling-transformation (CCT) diagram,

which is directly related to the chemical composition, prior

austenite grain size, and cooling rate. Figure 6 shows a typical

CCT diagram for SAE 1141-Nb (25) . Information from CCT diagrams is

essential to understanding transformation phenomenon, which in turn

affects the properties.

Probably the most effective way to increase a is to raise the C

content of the steel. An increase in C content from 0.1 to 0.6

mass % will increase the f(P) from <15% to about 90% for air

cooling of a large (200 mm diameter) bar. The effect of C content

on mechanical properties a-P steels is summarized in Fig. 7 (26)

The composition change will result in a nearly linear increase in

tensile YS, UTS, and transition temperature in notched impact

tests. The increase in C content will decrease both the reduction

of area in tensile testing and the upper shelf energy in impact

tests

10
Fig. 7. Effect of C content on mechanical properties of a-P

steels (25)

.

The most important structural feature that affects the strength of

P, for a given C content, is the spacing of the lamellar carbides

S . Usually the effect is described by a Hall-Petch type factor

where S

p

‘ 1/2 is the functional relationship. For eutectoid steel (C

content = 0.8 mass %) , evidence indicates that a S

p

_1 makes more

physical sense. Figure 8 shows data from Marder and uramfitt (27)

plotted as a function of S

p

_1

, which gives a positive value for o..

If the data were plotted as a function of S

p

/2

, o- for YS would be

negative, which makes no physical sense.

Lagneborg et al. (28) reported equations that relate the C and Mn

content to f(P) for steel bars with different diameters $, which

resulted in different cooling rates.

For 200 mm $ (dT/dt = 0.2°C/s in still air),

f(P) = 1.03* %C (1 + 0.74* [ %Mn ]

)

. (3)

For 70 mm $ (dT/dt = 0.5°C/s in still air),

f(P) = 1 . 18 • %C (1 + 0.64* [ %Mn] )

.

(4)

For 30 mm $ (dT/dt = 1.3°C/s in still air),

f(P) = 1.34* %C (1 + 0.60* [ %Mn] )

. (5)

11
Fig. 8. Effect of interlamellar spacing on YS of eutectoid

steel (26)

Depending upon the exact application, it may be necessary to also

consider other transformation products like bainite or martensite.

Bainite has a higher hardness and strength than a comparable a-P

mixture, but the toughness of bainite can vary quite a bit

depending upon the C content and type of bainite. Ollilainen et

al., (29) have shown that the lower temperature transformation

products in a a-P steel with V, Cr, and Mn additions and a C

content of 0.45 mass % increase strength, but lower the toughness

and machinability . For a steel with a lower C content (0.25 to

0.35%), bainite has been shown to improve both strength and

toughness (13,30). Given the variation in section size of many

forgings, like that shown in Fig. 9, the cooling rate in still air

after forging could result in bainite and/or martensite formation

in the thinner sections. The CCT diagram is useful in this regard

because it will define the cooling rates for which bainite-martensite formation can occur.

12
Fig. 9. Schematic diagram of cross sections through two typical

forgings. The effective cooling rates for specimens T8 and T9 will

be lower than that for specimen T1 due to the different section

sizes for the two forgings.

1.4. 1.4 Effect of Solid Solution Strengthening

Solid solution strengthening is not as important to YS as grain

size, transformation (pearlite content) and precipitation, and is

probably more important in terms of hardenability effects on the

CCT diagram. Normally ass is added to the other terms as a constant

times the content of a given element (mass %) . The value of the

constant is determined empirically, but is related in theory to the

lattice distortion that the element produces in the ferrite matrix.

Mn and Si are the two most important substitutional alloying

elements. The only interstitial element that normally needs to be

considered is N. In theory, the Mn and Si can be assumed to be

"free" and in solution (in most cases, part of the Mn and Si is

actually tied-up as inclusions) , but the "free" N content is

generally assumed to be zero when MA elements are present.

1.4. 1.5 Effect of Texture and Dislocations

These two factors are relatively unimportant for the type of

forging steels discussed here.

1.4. 1.6 Summary of Strength Effects

Gladman and Pickering (18) have summarized the results of numerous

tests on a-P steels without MA and interpreted the data in terms of

the effects of different strengthening mechanisms. Statistical,

13
multiple-regression analysis of the data produced the equations

shown below that predict YS reasonably well (±48 MPa)

YS = f (a) 1/3 • { 35 + 58 (mass % Mn) + 17.4*d' 1/2

}

+ [1-f (a) 173

]

• {178 + 3.8*

* 1/2

} . (6)

P and a slightly simpler

YS = f (a) 1/3

• { 35 + 17.4* d' 1/2

) + [1-f (a) 1/3

] -178 + 3.8*

p

’ 1/2

, (7)

where f(a) is the volume fraction of polygonal ferrite, S

p

is the

interlameliar spacing of the pearlite and the microstructural

features are measured in mm. YS varies nonlinearly with f(a) and

has been described by a modified rule of mixtures involving an

exponent of 1/3 on the volume fraction of proeutectoid a phase.

The factors controlling YS of the individual components are on

different scales. For a, YS is largely a function of the grain

size, which is typically in the 5 to 50 /Ltm range. For P, YS is a

function of the inverse square root of S

p

, which can vary from

about 0.1 to 0.4 /m.

Burnett (31) used a regression analysis to describe YS of MA a-P

steels by forging parameters and chemical composition over a

limited range of variables. The equation is shown below:

YS = 87.7 + 170 (%Mn) + 58.3 (Ln (dT/dt) ) , (8)

where dT/dt is in °C/min and % Mn is in mass %. There was

excellent agreement between measured and calculated values (0.99 R~

square and 95% confidence band is ±2 0 MPa) , but it does not address

the fundamental factors (<r , a and o .) that contribute to YS.

There have been many attempts to relate a Hall-Petch type relationship between grain size and either hardness or flow stresses at

higher strains (32)

Such relationships have to be viewed with

skepticism from a theoretical standpoint because the strainhardening effects are complex and difficult to model except for

specific examples.

For the UTS of a-P steels without MA, Gladman and Pickering (18)

report an equation, similar to Eq. (7) for YS, that is shown below:

UTS (±50 MPa) = f(a) 1/3 {246 + 18.2-d' 1/2

}

+ [ 1-f (a) 1/3

j (720 + 3.5*S

p

" 1/2 }. (9)

Equation (9) describes 96% of the data.

14
Lagneborg et al. (28) have pointed out that the UTS is more

important than YS for evaluation of MA a-P steels in automotive

applications because it relates better to component performance,

like fatigue resistance. Therefore, they propose an equation,

UTS = 177 + 533 • { f (P) + (2.2 + 0 . 71 • log10 [dT/dt ] )

•

(5 *N + V) + 0 . 12 • Si } , (10)

where f(P) is the volume fraction of P, N, Si, and V are in mass %

and dT/dt is the cooling rate at 750 °C in ° C/sec. The terms for

cooling rate, N, and V are related by the kinetics of transformation directly to the size and volume fraction of precipitation.

Even though Eq. (10) for UTS has no fundamental relationship to

theory, it certainly is easy to use because it incorporates precipitation effects through the cooling rate and composition terms,

rather than in terms of structural features. The comment about the

significance of YS to the application is well taken. Equation (10)

does not incorporate any terms related to S , probably because the

variation in S

p

normally found in air-cooled steels has relatively

little effect on UTS. The equation also allows for tradeoffs

between strengthening mechanisms, e.g.

,

reduced P content can be

made up by increased cooling rates and/or higher V and N contents,

which produce more precipitation strengthening.

Several potential problems are inherent with the statistical,

multiple-regression analysis of test data, which considers the

strengthening mechanisms as additive, like those published by

Gladman and Pickering (18) and Lagneborg et al. (28) . Often there

is doubt as to whether truly independent variables have been

considered in the analysis. The size scales over which the

strengthening mechanisms operate may dictate that one mechanism

will dominate over the others.

1.4.2 Hardness

Hardness is dependent upon P content and precipitation because it

is really associated with plastic flow just like the ultimate

strength. Standard Rockwell tests with C or B scales and Brinell

tests are macroscale hardness measurements because the size of the

indention and the locally deformed area are generally greater than

the size of the microstructural features. The hardness is an

average over all of the features. Vickers and Knoop hardness tests

are microhardness scales because the size of the indention is

similar to the size of the microstructural features.

Correlation between hardness and tensile properties can be made for

specific grades, like SAE 1141. These types of correlations are

15
extremely valuable when comparing steels with similar composition,

but trouble is expected when the correlations are extended to

different alloys or the same alloy but with different microstructures . 1.4.3 Toughness

After strength, the major issue related to forging steels is

toughness. The physical phenomena related to toughness are more

complex than that considered for strength. At lower temperatures

and higher strain rates, resistance to cleavage is important. Near

ambient temperatures and at slow strain rates, microvoid coalescence is important. For cyclic loading, the formation of fatigue

striations leads to failure.

1.4. 3.1 Impact Testing

The most frequently used toughness evaluation in industry is the

CVN impact test. Figure 3 showed the results of CVN testing from

different types of steel over the range of temperatures where the

fracture mode changes from ductile to brittle. For forging steels,

the common measure of toughness is the ductile-brittle (cleavage)

transition temperature (DBTT) . There are several ways to define

DBTT, but the most common way and the way used here is the

temperature at which the energy absorbed during the test is 27 J

(20 ft lb) . In general, the DBTT as well as the upper and lower

shelf energies may be important.

Of the three parameters that can be used to characterize CVN

testing (DBTT, upper shelf energy, and lower shelf energy) , microstructure has the biggest effect on DBTT, and as a result, DBTT has

been studied extensively (33-35)

Again, similar to the empirical

modeling that has been done for strength, grain size is an extremely important factor affecting DBTT. The effect of a grain size, or

the P colony size p, on DBTT is due to an increasing difficulty in

propagating a cleavage crack as the grain size decreases. Each

time a crack encounters a high-angle grain boundary, the crack has

to re-orient itself. Experimental evidence has shown that grain

boundaries can stop a crack so that it must re-nucleate at a second

location. The value of p is determined by the prior austenite

grain size D and to a lesser extent by the cooling rate.

The energy absorbed in a CVN test at room temperature for a MA a-P

steel is generally in the transition region (see Fig. 3) (15) . A

higher resistance to ductile fracture, or a lower resistance to

cleavage fracture will change the transition region and could

influence the results. At a given value of austenite grain size,

toughness is increased by a lower C content (more a, less P) , and

more P dilution (increasing the cooling rate and Mn content) , but

these effects are secondary in comparison to that of D and p.

Given the complex interactions between microstructural

constituents, CVN testing at room temperature may not be a

16
clear-cut way to characterize impact toughness. The whole

transition curve is valuable for a better evaluation.

The energy absorbed during impact testing at temperatures below

DBTT, referred to as the lower shelf energy, is related to cleavage

fracture on {100} planes in ferrite. Microstructural effects on

the lower shelf energy are of little interest in the technical

community because the absolute value of the absorbed energy is low

(<5 J) , so that even large relative effects are insignificant.

The energy absorbed during impact testing at temperatures above

DBTT, referred to as the upper shelf energy, is related to ductile

fracture by microvoid coalescence at second-phase particles in the

structure and can vary significantly with microstructure. The most

important factors to consider are the volume fraction of second

phase and strength, which are largely determined by the f(P) and

precipitation. Figure 7 (26) showed how increases in C content,

which determines the f (P) , improves the strength, but lowers toughness (reduction in area (RA) goes down, DBTT goes up, and the upper

shelf energy is lower)

.

The effects of structure on DBTT for a-P steels have been described

by a simple rule of mixtures that incorporates each constituent.

Gladman and Pickering (18) give the following equations to predict

the DBTT:

DBTT (±30 °C) = f(a)*[-46 - 11.5*d' 1/2

] + (1 - f(a))-[-335

+ 5.6* S

* 1/2 - 13.3-p' 1/2 + (3.48X10*6

) -t] , (11)

r

that explains 82% of the data. There are complex interactions for

the microstructural constituents, i.e., p, d, S

p

, f(a), and carbide

thickness t, which are not apparent from Eq. (11) . Decreasing t

requires S

p

to increase, so that the rate of change in DBTT is not

independent of the two factors. For a given C content, the

relationship between t and S can be described by the following

equation (36,37):

t = S

p / {[f(P) / 0.15 (% C)] -1} . (12)

The influence of p, f(P), and d on DBTT are illustrated in Fig. 10.

Higher f(P) increases DBTT by reducing the influence of a grain

size in the first term in Eq. (11) . It is important to note that

the effect of a grain size on the CVN test is large even in the

absence of P, see Fig. 10 where the DBTT at zero f(P) for the two

initial grain sizes differ by nearly 50 °C.

17
Fig. 10. Interaction between pearlite fraction, pearlite colony

size i and ferrite grain (/m) size on the DBTT in a-P steels

(18) .

Strength has a large effect on DBTT, but is not specified in Eq.

(11) because strength depends on the same structural factors that

influence DBTT. If the equations for YS and DBTT are considered

together, then the DBTT will increase by about 0.3°C/MPa increase

in YS.

Recent developements in Japan have shown that impact toughness can

be improved by intragranular a formation in free machining (high S)

steels that also contain V (38) . Figure 11 shows the energy

aborbed in a Charpy U-notch impact test as a function of temperature for three different steels: air-cooled 0 . 25C-1 . 5Mn-0 . 07S-0 . IV

steel, air-cooled 0 . 25C-1 . 5Mn-0 . 02S-0 . IV steel, and a Q&T 0.45C0.8Mn-0.02S steel. The air-cooled, high S, V steel is significantly better than the low S version, and slightly better than the

Q&T steel at all temperatures. It is important to note that the

results from a U-notch test can be different from a standard CVN

test due to the lower stress concentration at the notch (compare

the differences between the results shown in Figs. 3 and 11)

In the USA, Kirby et al. (37) have shown a lower DBTT (better

toughness) in CVN tests for a high S, V, 0.4C steel compared to a

similar steel without high S or V, but the trend does not appear as

significant as suggested by U-notch results shown in Fig. 11. The

mechanism responsible for the increase in toughness is still being

studied, but it appears that the high S and V content changes the

transformation kinetics for proeutectoid a. The higher toughness

is related to grain refinement and increased f(a), and could be

described by an equation similar to Eq. (10)

18
Fig. 11. U-notch impact toughness of steels as a function of test

temperature (35)

.

1.4. 3.

2

Fracture Toughness Testing

Fracture toughness (Klc ) testing is a well-known technique used to

evaluate toughness of materials, and was developed as a tool to

determine the f itness-for-purpose of a particular material and

processing, given the service operating stress and existing flaw

size. However, KIc testing has not been used routinely for characterization of materials, except for critical safety applications.

KIc testing is expensive, and it takes a relatively large specimen

to get good results.

Limited KIc testing has been used to compare MA a-P to Q&T steels.

Ollilaninen et al. (29) found that KIc was four times higher for Q&T

steels compared to MA a-P steel at a YS of 650 MPa. The difference

in KIc between the two steels was not considered significant for

many applications because KIc did not change the fatigue life, the

limiting failure mode, considerably. A schematic diagram

illustrating the fatigue behavior for Q&T vs. MA a-P steels is

shown in Fig. 12. Given the fact that most of the fatigue life

19
Fig. 12. A schematic curve of the fatigue crack growth and final

fracture of high-strength MA a-P steel and Q&T steel (29)

is consumed with initiation of crack growth, the difference in

number of cycles to failure is rather small compared to the

difference measured in K. . IC Efforts to relate KIc from other mechanical properties have been

attempted ever since fracture mechanics became accepted in the

technical community. Figure 13 shows KIc for different MA a-P

steels plotted against the UTS for different steels (9) . The

correlation between KIc is not encouraging. There does appear to

be a minimum value of KIc of about 50 MPa*m1/2 for these types of

steels.

20
KSI

105 115 125 135 146 155

1

S.

z

8

Ui

Z

X

o

=>o

H“

UI

DC3

H*O

<

oc

u.o

700 800 900 1000 1100

ULTIMATE TENSILE STRENGTH, MPa

Fig. 13. Correlation between fracture toughness and UTS for a-P

steels (9) .

Short-rod, chevron-notch fracture toughness (KIv ) testing is a

relatively new development in fracture mechanics that has several

advantages compared to standard fracture toughness testing. First,

the specimen is 2.5x smaller than that required for compact tension

testing of ASTM E 399 plane strain fracture toughness. There is no

fatigue precracking required for the test as in other fracture

toughness test procedures. Past experience (30,39,40) has shown

that KIv has excellent potential for use as a screening test to

evaluate high-strength aluminum alloys and steels. Figure 14 shows

a excellent correlation between K

lc

and Klv for Q&T steels at room

temperature

.

1.4. 3.

3

Tensile Testing

At present, little notice is given to the fracture event of a

tensile specimen? typically, the actual fracture load is not even

recorded for a test. Of the commonly reported information from a

tensile test, only the percent reduction in area (RA) and to a

lesser degree, elongation are influenced by fracture. Neither of

21
Fig. 14. Correlation between fracture toughness of Q&T steels

determined with a compact tension and a short rod specimen (9)

these properties has received much attention in the literature

compared to the extensive work done on YS and UTS.

Figure 15 shows the effect of second-phase particles, volume

fraction and shape, on fracture strain (23) . The true strain to

fracture e

T and %RA decrease with an increase in volume fraction of

second-phase particles. Spherical particles are less detrimental

to ductility than elongated or plate-like particles.

Figure 16 shows e

T plotted against the upper shelf CVN energy (18) . A good correlation is found between tensile and impact test

results. Figure 17 shows the true strain plotted vs. f(P). Most

of the strain is concentrated in the neck as nonuniform strain.

There is an exponential decrease in both the uniform strain e

u and

e

T with increasing f(P).

Pickering (41) reported the following equation to predict the

tensile ductility from the microstructure of a-P steels:

%RA = 78.5 + 5 . 39 ( %Mn) - 0.53(% P) - 0.328(d) , (13)

22
Fig. 15. Effect of second-phase particles and their morphology on

total ductility in uniaxial tension (22)

.

Fig. 16. Plot of true fracture strain in uniaxial tension vs

upper shelf energy absorbed in CVN test (18).

23
Fig. 17. Effect of pearlite fraction on the uniform, e

u

, and

fracture strain, e

T

, in uniaxial tension for a-P steels (18)

where d is the mean linear intercept grain size in micrometers

From the few reports in the literature, grain size seems to have a

large effect on %RA. An order of magnitude decrease in grain size

from 100 to 10 /im increases the %RA by 30%. Ollilaninen (29) shows

a similar dependence of %RA on grain size for MA a-P steels.

In terms of the fracture stress ofr , the only data available in the

literature are for eutectoid steels with 100% pearlitic microstructures. Gladman and Pickering (18) report that afr is related

to S

p

by a Hall-Petch relationship

afr = 790 + 8.5-S

p

' 1/2

. (14)

The relationship is interesting because the constant term in the

Hall-Petch relationship, cr . , increases from 178 MPa for YS, to 550

MPa for e = 0.15, and then to 790 MPa for fracture. The coefficient for S goes from 3.8 at YS, 10 at € = 0.15, and 8.5 for

fracture. The trend is towards higher values of a j and k

y

in Eq.

(1) for increasing strain.

Marder and Bramfitt (27) found a similar relationship between S

p and cr fr in a tensile test of 100% pearlitic steel, but with a

different exponent:

afr = 436 + 9.81-Sp' 1

, (15)

24
where stress is in MPa and spacing is in ran. Again, in this case,

a better fit to the data was obtained for S

p

_1 compared to S

p

' 1/2

. The optimum microstructure needed to maximize af has not been

determined. The Fe

3

C constituent in P is the main contributor to

the volume fraction of second phase particles in MA forging steels,

so we know that we need to minimize the carbide phase in order to

obtain the highest possible %RA (Fig. 15). But a high f(P) is

needed to raise the flow stresses (Fig. 4) . There should be an

optimum P configuration, t and S

p

, for a given C content that optimizes cr fr and %RA.

A modified Bridgman analysis of the stresses found in a tensile

specimen has been proposed as a means of characterizing the toughness of a tensile specimen (42)

.

The sum of the equivalent plastic

flow stress o

Q

and hydrostatic stress aT represents an approximation

of the critical stress oc required to nucleate the majority of

voids that are found on the fracture surface. Figure 18 shows the

results of the analysis on a series of tests at different temperatures run on an austenitic stainless steel with two processing

histories. The analysis allows the degree of sensitization to be

quantified in a simple manner. The critical stress for void

nucleation can be used as a failure criterion, one which is linked

to the physical process related to fracture, for stress analysis of

components in more complex stress states (43)

.

TEMPERATURE, K

Fig. 18. Plot of interfacial stress at fracture in uniaxial

tension vs. test temperature for two austenitic stainless steels,

annealed and sensitized (42)

.

25
To summarize the effects of microstructure on fracture in uniaxial

tension, there is limited evidence that fracture properties are

sensitive indicators of microstructural features, just like YS and

UTS. In the future, correlations between tensile fracture

properties vs. fracture toughness and fatigue resistance are

needed. With additional experience, tensile fracture properties

could be more valuable.

1.4. 3. 4 Fatigue Testing

Fatigue is the typical failure mode for many forged, automotive

components, but fatigue testing is not the main procedure for

characterizing materials and processing. The reason that fatigue

testing is not more prevalent to characterize materials and

processing effects is because fatigue testing is expensive and time

consuming, and the results are sensitive to the applied stress

level. The stress concentration due to the design of a component

and the tensile properties are both critical factors that determine

fatigue life. Designers usually depend upon correlations between

tensile properties or hardness and fatigue life to predict a

material's fatigue life in the preliminary stages. Then fatigue

testing is often performed in the final stages of development on

the actual components after the material and processing schedule

decisions have been made.

Hardness and UTS are the most important factors related to the

endurance limit for fatigue (the stress range required for a life

of 106 load cycles) . Figure 19 shows the correlation between the

fatigue endurance limit and hardness. Hardness and UTS are

influenced mainly by precipitation and f (P)

,

which in turn are controlled by chemistry and cooling rate. Typically, the endurance

limit for a MA a-P steel is no more than 10% less than that for a

Q&T steel with the same hardness.

The behavior of MA a-P steels in fatigue has been studied in

strain-controlled fatigue tests as well as stress-controlled tests

for the endurance limit (9)

A comparison of typical cyclic and

monotonic stress-strain curves for MA a-P steel and Q&T steel is

shown in Fig. 20. The Q&T steel exhibits cyclic softening while

the stress-strain curve for the MA a-P steel changes little with

cyclic deformation. In applications where there is cyclic loading,

cyclic stability compensates for a lower monotonic stress-strain

curve. Figure 21 shows typical strain-life fatigue curves for a MA

a-P steel and a Q&T steel . The curves for the two steels are

similar, although the MA steel was slightly harder.

26
1000

Fig. 19.

hardness

Fig. 20.

steels (

20 30 40 50 60

ROCKWELL C HARDNESS

Correlation between fatigue endurance limit and

(19) .

STRAIN

A comparison of the stress-strain response of a-P and Q&T

)

•

27
Fig. 21. A comparison of the strain-life fatigue curves of a-P and

Q&T steels (9)

1 . 5 Summary

The effect of different strengthening and toughening mechanisms on

the YS and DBTT of steel is summarized in Fig. 22 (2) . Grain

refinement is the only strengthening mechanism that also improves

DBTT. All other strengthening mechanisms reduce DBTT. The effects

of strengthening mechanisms on other measures of toughness are not

quantified in the literature. For instance, grain refinement has

been known to decrease fracture toughness in Q&T AISI 4340 steel

(44) . Lagneborg et al. (28) sum up the current state of alloy and

processing design. The first goal is to maximize the strength,

which is used to design a component, without sacrificing toughness,

which generally determines the life of the component. Increasing

the P content by raising the C content will raise strength, but

decrease DBTT. It is possible to increase f(P) and keep the C

content constant by decreasing the transformation temperature,

28
P*53 N 2 *30

Fig. 22. Strengthening mechanisms and their effect on the DBTT

(I.T.T.) in steels with a predominantly polygonal a microstructure.

Ratios indicate the change in DBTT for a 15 MPa increase in yield

stress (2)

.

which is accomplished by increasing cooling rate and/or alloying.

In this way, the maximum strength for a given C content is

obtained. As the cooling rate and/or alloying increases at a

constant C content, S

p

decreases and a dilute P, one in which the

Fe

3 C platelette thickness is reduced, is formed.

Strengthening by the mechanism described above is limited by the

formation of lower-temperature transformation products like bainite

and martensite. The addition of V is an efficient way to add extra

strength to a-P steels, but again the higher strength comes with a

lower toughness.

The final goal of alloy design is to maximize the impact toughness

by minimizing the prior austenite grain size. Alloying with Al,

Ti, and N and close control of the forging process can keep the

prior austenite grain size relatively low. A problem arises

because no one can specify exactly how much impact toughness is

needed for a particular application.

29
The problem with low impact properties in a-P steels is somewhat

academic because fatigue is normally the limiting failure mode. In

application, there is no design requirement for impact toughness.

The high strain rate and notch used in CVN testing produces a

highly constrained specimen for evaluation that is not typical for

automotive applications. The CVN test is used frequently for many

reasons. It is a well-accepted procedure around the world, it is

inexpensive, and it has historical significance. The CVN test has

been used since World War II for quality control purposes. It is

extremely effective in identifying bad batches of low C and Q&T

steels. However, impact properties are of no relevance in relation

to the operating conditions for many automotive components and,

therefore, the same criteria should not necessarily restrict the

use of a new type of material, such as MA a-P steel (22). These

sentiments were repeated in private communications with several

researchers who work for steel companies and steel users (45-47)

It seems reasonable to conclude that impact properties are not of

prime importance in components of this type, which normally fail by

fatigue

Overall, KJc or KIv testing appears superior to CVN testing because

in fracture toughness testing, there is a direct link between

results and f itness-for-purpose. Problems remain in applying

fracture toughness testing on a wider scale for characterization of

materials. First, there is the question of acceptance. Many steel

users do not have the requisite expertise to apply fracture

toughness technology. Another factor could be financial, although

the use of short-rod testing would seem to minimize cost.

Characterization of fracture in tensile testing represents an

opportunity to supplement our understanding of toughness. Tensile

testing will have to be done in any case because tensile properties

are the main issue. The additional effort required to characterize

fracture in tensile testing is relatively small.

2. PROPOSED RESEARCH PROGRAM AT NIST

The program at NIST will address the research of structure-property

relationships in four areas. The first part is to generate

microstructures that are representative of thermo-mechanical

processing on the material of interest. The second part is

characterization of the pertinent structural features. The next

phase of research addresses the determination of properties which

are relevant to the potential applications for the material. The

final part of the program integrates the research results into a

model to predict properties from structure based on physical

metallurgy principles.

Figure 23 shows a schematic diagram that is used to describe

microstructural engineering research into TMP of steel . The end

result of a study of structure-property relationships will be a

30
Forging

Blank

Fig. 23. Schematic diagram of TMP modeling showing the

relationship between the structure-properties modeling and other

efforts

.

quantitative procedure that will take the significant structural

factors (output of transformation model) as input values and

predict their effect on properties, as related to performance as

much as possible. The ultimate goal is to be able to predict the

processing variables that will produce acceptable performance. In

the case described here for forging steels, the end product of

microstructural engineering will be a set of processing conditions

that will provide the desired fatigue life.

2 . 1 Generation of Representative Microstructures

A high-strain-rate servo-hydraulic mechanical test system with a

500 kN load capacity has been designed to simulate forging,

typically used in manufacturing of automotive parts and other semifinished bar products. Average plastic strain rates of about 10 to

15 per second are possible. A schematic diagram is shown in Fig.

24 to illustrate the loading apparatus to be used in the initial

testing. The specimens are solution heat treated in a convection

box furnace, using no special atmospherical control, and then

placed in the load frame. The hot specimen is supported on the

31
25 mm

Insulated

Stage

Base

Fig. 24. Schematic diagram of loading fixture for large TMP

simulator.

ends by spring-loaded, insulated columns. The specimen is then

given a prescribed deformation with a known strain rate. An

example of the load vs. stroke data for one experiment is shown in

Fig. 25. A sheet of graphite, 0.25 mm thick, is used between the

specimen and the dies to minimize friction during deformation.

Experimental variables that are important in generation of the

represenative microstructures for the initial survey of properties

are:

(1) reheating temperature before forging,

(2) strain during forging,

(3) cooling rate from forging temperature, and

(4) alloy composition.

32
Fig. 25. Load (vertical axis for the bottom line on the figure)

and stroke (vertical axis for the top line on the figure) plotted

against time (horizontal axis for both curves) data taken from TMP

experiment on SAE 1141 steel.

Figure 26 shows a schematic diagram that illustrates the

thermomechanical variables to be used in the first series of tests.

The initial reheat temperature will be either 950, 1100, or 1215 °C.

Typically, the reheat temperature for forging is 1200 to 1300 °C,

allowing V and Nb to dissolve in the austenitic matrix. After

reheating, the specimens will be furnace-cooled to either 1100 or

950 °C before deformation so that the finish-forging temperature is

close to 1100 or 950 °C, respectively. The deformation associated

with forging is usually finished with minimal heat losses (finishforging temperature >1100 °C) , making plastic flow easier by

reducing the loads required for forging and minimizing die wear.

Some forging shops are working with lower reheat and finish-forging

temperatures in order to save on energy costs and to limit decarburization that must be machined off later.

A high and a low degree of deformation (e of about 0.5 and 0.2,

respectively) will be done for selected conditions. After deformation, specimens will be cooled to RT at different rates, so that

the effect of cooling different section sizes in still air can be

evaluated.

33
oouJ

£E3

K

<

CC

UiQ. 5

UJ

8=

o

UJ

CC3H

<

DCUJ

a.5

UJ

Fig. 26. Schematic diagram showing the four different TMP

schedules to be evaluated in the initial study.

For the initial phase of our experimental program, we will use two

different alloys, SAE 114 l-V and -Nb. The composition of the

alloys is shown in Table 2. Both are classified as free-machining

grades due to the high S content. The main difference between the

two steels is the MA additions of V in one case and Nb in the

second. Within the next year, we expect to obtain additional

steels with slightly different alloying additions so that the

effects of the high S content and MA can be studied further.

Table 2. Chemical composition (mass %) of steels for test

program.

c Mn P S Si Cr N MA element

0.37 1.44 0.012 0.100 0.25 0.08 0.009 0.07 V

0.42 1.58 0.013 0.100 0.25 0.09 0.009 0.026 Nb

34
2.2 Characterization of Structure

The characterization will include standard metallographic

procedures for light and electron microscopes. For the study of MA

a-P steel, the relevant features to be measured include f(a), S , a grain size, p, D, f(ppt), and x t

. For the initial survey of

properties, transmission electron microscopy (TEM) examination is

not feasible for every processing condition, but can be evaluated

later for selected conditions. In addition to microstructural

features, it could be important to document the microhardness (DPH)

of each phase (a and P) where possible. DPH could be valuable for

two reasons. First, where S varies and is discontinuous, DPH

could give a direct estimate of the strength of the P phase without

measuring S . The characterization of the P colonies is important for several

reasons. First, f(P) is a big contributor to the strength and is

related to different factors such as S

p

, t, and p. The scale of

the features is small, S

p

can be as fine as 65 nm (48). TEM

studies may be required to accurately describe the features of P.

Thermal analysis of the cooling curves after processing could help

characterize the features because S

p

is mainly a function of the

temperature for transformation (49)

.

Additional study of selected broken specimens will be done to

document the fracture mechanisms. SEM observation of fracture

surfaces and polished sections through failed specimens can be used

to associate damage mechanisms to microstructural features (this

part has received little attention in the literature.

2 . 3 Determination of Properties

Standard tension testing of cylindrical specimens and hardness

testing (Rockwell C or B scale) will be the standard procedure for

experimental evaluation of strength. In the initial survey of

properties, %RA and the stresses at fracture will be used as

measures of toughness. The procedure for calculating oo and aT are

found in Appendix A. For selected processing conditions, chevronnotch fracture toughness tests at RT, CVN transition curves, and

possibly fatigue testing can be included for a more complete

evaluation of toughness.

Machinability is also an important consideration in the choice of

material for many applications (4,13). Typically, the evaluation

of machinability is determined by measurement of tool life during

different operations like milling, turning, and hole drilling,

requiring a large volume of material which is not available in a

small-scale simulation. Machinability is generally assumed to be

a function of hardness, a lower hardness is easier to machine. In

practice, there is a range of hardness specified for a given

application so that the material will be strong enough to withstand

the loading during application, but soft enough to be machined.

35
Empirical correlations between hardness and machinability , and

hardness and tensile properties are used to predict performance.

2.4 Integration of Structure and Properties for Modeling the

Application

The prediction of mechanical properties from structural features is

the goal of our research. Considering the prediction for YS and

UTS of a-P steels, the existing equations from Gladman and

Pickering (18) are adequate to evaluate strength. For the

precipitation strengthening contribution, Eq. (2) is the best way

to quantify the effect; it is well supported by theory and limited

experimental evidence, but requires extensive TEM observations for

application. Simpler experimental procedures for the

characterization of fine-scale precipitation would be extremely

helpful in this regard.

For the prediction of toughness, tensile fracture characterization

is key. A test procedure to evaluate toughness other than CVN

testing, particularly one that includes a f itness-for-service

analysis, could aid significantly in application of MA a-P steels

to a wider range of use. By characterization of the stresses at

fracture and calculation of oc for the initiation of voids, toughness is related to the physical process of ductile fracture and

distinct microstructural features. The interpretation of oc for

void nucleation from microstructural features is new and the

relevance is open to debate. However, the procedure for the

characterization of the stresses at fracture is simple and easily

applied. Correlations of oc vs. K

Jv

and oc vs. fatigue life could

be studied and would provide a basis for evaluation of the tensile

fracture characterization.

3. PRELIMINARY STUDY

3.1 Introduction

In this part of the report, we will show preliminary results from

our initial study on the thermomechanical processing (TMP) of MA

a-P steel. The goals for this study are as follows: (1) to

illustrate the facilities that are available at NIST for the study

of the TMP of steel, (2) to compare the results of the larger

simulator to previous results from earlier work done here at NIST

on smaller specimens, and (3) to evaluate the existing equations

that predict properties from microstructure.

3 . 2 Material and Experimental Procedures

The composition of the steel used in this preliminary work, shown

in Table 3, is similar to SAE 1141 with a Nb addition. This alloy

has been used is several recent studies by Cheng and co-workers

36
(25,50,51). The material was received in the form of a 30-mmdiameter bar,' and the hardness was about R 18. c

Table 3. Chemical composition of SAE 1141 with Nb (mass %)

.

c Mn P S ; Si CU Ni Cr Mo N Nb

0.42 1.49 0.015 0.099 0.26 0.16 0.06 0.06 0.009 0.0075 0.039

The microstructure of the as-received bar in the transverse

orientation is shown in Fig. 27. The pearl itic phase occupies most

of the area while the a phase, 5 to 8% of the area, is present as

a network at the prior austenite grain boundaries. The prior

austenite grains vary in size, but the average diameter is about 40

jum. Part of the reason for the variation in prior austenite grain

size could be the nonuniform distribution of MnS particles in the

structure. At the highest magnification (Fig. 27b) , spacing of

carbide platelettes in the pearl it ic phase can be observed with the

light microscope in some areas. The finest spacings could only be

observed in SEM. The minimum S

p

observed was about 0.15 /m. The

true spacing should be equal to the finest observable spacing.

In Fig. 28, the MnS inclusions are shown in longitudinal and

transverse sections. In the transverse section, Fig. 28a, the

particles are largely spherical and located in the a network that

formed at the prior austenite grain boundaries. The particles

appear to have pinned the austenite grains during processing,

preventing grain growth, as suggested by others (13) . When viewed

in the longitudinal section. Fig. 28b, the inclusions are strung

out along the rolling direction.

Mechanical tests were conducted on the as-received material as well

as the processed steel so that the effect of subsequent TMP could

be quantified. Hardness testing was done on the Rockwell C scale

(Rc ) to determine the bulk hardness, and on the Vickers scale (DPH)

with 100 and 10 g loads to determine the micro-hardness. The size

of indention in the DPH test is an important consideration. For

the 10 g load, the indention is correspondingly small, but the

scatter in results may increase as the load decreases. The 100 g

load should produce more uniform results, but is less discriminating with regard to the individual microstructural features,

particularly, the hardness of a in a thin network may be difficult

to characterize.

Tensile and CVN specimens were oriented along the rolling direction

of the bars. The tensile specimens were cylindrical in shape with

a 6.4 mm diameter and a 25 mm long gage section. The crosshead

rate for tension testing was 0.5 mm/min. In addition to the

typical tensile properties that are reported, values for o

Q

37
(a)

50 /m

Fig. 27. As-received microstructure of SAE 1141-Nb steel (a) low

magnification, (b) high magnification.

38
Fig. 28. Appearance of MnS particles in microstructure:

(a) transverse section and (b) longitudinal section.

39
(equivalent plastic flow stress at fracture) and oc (critical

stress for void nucleation at fracture) (42) are described in

Appendix A. CVN specimens were machined according to ASTM E 23,

type A.

Specimens were processed in TMP simulator described earlier. The

specimens were deformed to a true strain of about 0.8 at an average

strain rate of 15. The temperature varibles in the experiments are

listed in Table 4.

Table 4. Temperature variables used in preliminary study.

Spec.

Reheat

Temperature

Deformation

Temperature

Cooling

Rate*

Dl * * *

D2 900° C 900° C 0 . 8°C/s

D3 1100°C 1100°C 0 . 9°C/s

+ Cooling rate is defined by the time to cool from 750 to 500° C. * As-received condition.

3 . 3 Results

3.3.1 Microstructure

The microstructures generated during TMP are shown in Figs. 29 and

30. For the D2 condition (900°C deformation, shown in Fig. 29),

the a-P features are much finer than that of D1 (as-received) . The

a phase occupies between 35 and 43% of the cross sectional area.

The a grain size is about 5 /xm while p is slightly larger, about 10

/xm average diameter. S

p

for D2 is variable, but in general appears

finer than that found in Dl, about 0.13 jixm.

For the higher processing temperature (1100 °C, condition D3)

,

the

microstructure shown in Fig. 3 0 is similar to that of Dl (asreceived) , but refined slightly. The prior austenite grain size is

about 2 0 jLtm, and the a content is about 8 to 10% of the area. The

value of S

p

for D3 is similar to that for Dl, about 0.15 /xm.

3.3.2 Mechanical Properties

The mechanical properties at RT are summarized in Tables 5 and 6.

The tensile properties do not change significantly with the

different processing conditions. The CVN test results are

interesting because condition D2 failed by a fully ductile mode.

The energy absorbed in testing D2 was about 50% higher than the

40
Fig.

1100

29. Microstructure of SAE 1141-Nb steel after deformation at

C and air-cooling (a) low magnification, (b) low magnification.

41
Fig. 30. Microstructure of SAE 1141-Nb steel after deformation at

900 °

and air-cooling (a) low magnification, (b) high magnification.

42
average for conditions D1 and D3, where the failure mode was a

mixed ductile-brittle one. The properties are typical of a a-P

steel (compare the data for 1141-Nb with that shown in Fig. 2)

.

The hardness testing is summarized in Table 7. Given the size of

the indention and the thin network of a, DPH of a could not be

determined in some cases. The hardnesses of D1 and D3 are similar,

as were the microstructures and mechanical properties. The bulk

hardness for condition D2 is lower than the others. For the

pearlitic phase, the lower DPH for the 100 g load is probably due

to the soft, ferritic matrix that surrounds the relatively small P

colonies. With the lighter load, the pearlitic phase in condition

D2 has the highest DPH because the indention is small enough to

only sample P, and not the ferritic phase. The average microhardness for a and P is always higher than the average bulk

hardness. The scatter in DPH testing is greater for the lighter

load.

Table 5. Tensile test results at RT on SAE 1141-Nb steel.

Specimen #

YS,

MPa

UTS,

MPa % Elong

.

% RA

Dl-1 475 760 15.6 30.0

Dl-2 467 748 15.1 29.3

D2 486 703 19.2 32.5

D3 493 774 16.3 30.0

Table 6. CVN test results at RT on SAE 1141-Nb steel.

Specimen #

% Cleavage

Fracture

Absorbed Energy,

J

Dl-1 48 33.5

Dl-2 61 22.2

Dl-3 42 37.3

Dl-4 50 26.0

D2 0 43.4

D3 38 25.2

43
Table 7. Summary of hardness testing.

Process

Condition Brine It — bulk DPH, 100 g load DPH, 10 g load

Dl -- a

221, 217, 217

214, 215

Dl -- P ave. =217 269,266,290,265,275

ave. = 273 (Brinell 264)

290,237,254,279,263

ave. = 265 (Brinell 253)

D2 -- a

203, 202, 202,

202, 203

ave. = 203

221,279,214,219,260

ave. = 239 (Brinell )

02 -- P 212,224,227,243,226

ave. = 226 (Brinell 219)

279,286,293,313,325

ave. = 299 (Brinell 286)

D3 -- a

221, 222, 222,

224, 222

ave. = 222

297,263,226,219,260

ave. = 253 (Brinell 243)

D3 -- P 258,252,264,262,243

ave. = 256 (Brinell 243)

245,297,263,297,272

ave. = 275 (Brinell 264)

The tensile fracture stresses are shown in Table 8. These

stresses, along with the %RA, characterize the ductile fracture

behavior of the material. The critical stress is essentially

constant for the three different processing conditions. Equation

12 can be used to predict the effect of changes in the microstructure on % RA. For condition Dl, the predicted % RA would be 32%,

agreeing well with the measured value of 30%. However, for

condition D2 , the prediction, 47%, is considerably higher than the

measured value, 33%.

Table 8. Tensile fracture stresses (MPa).

Specimen # *1

Dl—

1072 770 1842

Dl-2 1026 727 1753

D2 1129 686 1815

D3 1076 751 1827

3.3.3 Fractography

The appearance of the tensile specimens from all three conditions

was similar. The fracture surface from an as-received tensile

specimen is shown in Fig. 31. At 5Q0x, a few irregular troughs

44
Fig. 31. Fracture surface of tensile specimen taken from asreceived steel (a) low magnification, (b) high magnification.

45
and ridges, 50 to 100 /xm in size are visible. At higher magnification, dimples associated with inclusion particles (labeled A) are

the dominant features on the surface, but smaller facets (labeled

B) are found between patches of dimples. The small facets are

areas where the fracture went through a pearl ite colony.

For condition D2 (processed at 900 °C)

,

the tensile fracture surface

is shown in Fig. 32. The higher a content in the microstructure

has little influence on the appearance. The dominant features are

inclusion-nucleated dimples. Patches of submicrometer-sized facets

are found on the surface between the larger dimples. The consistent

tensile fracture appearance can be related to the similar fracture

stresses of the tensile specimens.

The appearance of CVN fracture surfaces is different from that of

the tensile specimens. For conditions D1 and D3 , large areas of

cleavage fracture are found on the surface (see Table 6) . For

condition D2 , no cleavage is observed on the fracture surface, but

the features are still diffrent from that observed on the tensile

fracture surface. Figure 33 shows a representative area from the

region directly ahead of the notch tip where the fracture process

started. Numerous dark cavities, 10 to 50 /xm in size, are prominent on the surface between the other fracture surface features.

The cavities are smaller, 5 to 10 /xm, than those on the tensile

fracture surface in Fig. 32. Otherwise, the fracture surfaces of

the tensile and CVN specimens are similar.

3.4 Discussion

The three different processing conditions considered in this

preliminary study have altered the microstructure; however, the

measured tensile properties are only slightly affected by the

change in features, the biggest difference is in YS between D1 and

D3 , about 10%. CVN energy absorbed at RT does significantly

increase for the finest ferrite grain size (condition D2)

There are considerable microstructural differences between

conditions D1 and D2 . The a content increases from about 5 to 40%

of the volume, and S

p

decreases from 0.15 to 0.13 /xm. The a grain

size for condition PD1 is taken as the thickness of the grainboundary network, about 3 /xm, while the a grain size is about 5 /xm

for condition D2 . It is helpful to consider the changes in

microstructure and their effect on tensile properties, related to

Eqs . (7) and (9) . Kirby et al. (36) used these same types of

equations and found excellent agreement between the measured

tensile properties of a a-P steel which was not precipitation

strengthened and predicted values. For V-containing steels, the

equations underpredicted the strengths.

46
Fig. 32. Fracture surface of tensile specimen taken from

deformed at 900 °C and air-cooled (a) low magnification, (b)

magnification.

steel

high

47
Fig. 33. Fracture surface of CVN specimen taken from steel

deformed at 900 °C and air-cooled.

Table 9 compares the predicted values for YS and UTS from Eqs. (7)

and (9) for conditions D1 and D2 with measured values from Table 5.

The predictions are 10 to 15% off the measured values, overpredicting the properties for Dl, and underpredicting YS and UTS

for D2 . In this case, the under prediction is not related to

precipitation strengthening because the Nb-treated steel used in

this study does not contain any V.

Table 9. Comparison between measured and predicted tensile

properties for SAE 1141-Nb steel.

Condition

Eq (7)

YS, MPa

Measured

YS, MPa

Eq. (9)

UTS. MPa

Measured

UTS, MPa

Dl 521 471 848 754

D2 405 486 641 703

One possible reason for the problem in predicting strength is the

characterization of the pearlitic phase. The measurements of S

p and f(P) are probably not sufficient for two reasons. First, the

48
pearlitic phase of the Nb-treated steel used in this initial study

has a discontinuous carbide morphology in many regions. Figure 34

shows a typical region where there are numerous P colonies along

with the ferrite network and MnS particles. The carbide constituents in the pearlitic phase do not always form the classical

fingerprint pattern that can be characterized by a single parameter, S

p

. The second factor is that the dilution of P depends upon

the processing conditions. For condition D1 and D3 , where there is

little proeutectoid a, the dilution is high. For conditon D2

,

where the a content is much higher, the dilution is relatively low.

Additional work is needed to improve the characterization of the

pearlitic features.

An important consideration of the test results is the correlation

between tensile strength and hardness. The correlation is

important because it relates directly to the allowable hardness

range that is specified in design of components. Figure 35 shows

YS and UTS for the steels tested here and other data from Kirby et

al. (37) on similar steels plotted against the bulk hardness

determined on the Rockwell C scale. The YS predicted for condition

D2 in this study from the trend observed for Kirby's data would be

less than the measured YS.

Fig. 34. Microstructure of steel in as-received condition showing

the difference in pearlite morphology in different colonies and the

MnS inclusions in the a network.

49
The evaluation of toughness can be made by comparing the results of

tensile fracture properties and the CVN results. The ratio a

pr

at

fracture/YS is taken as a measure of toughness from the tensile

test results. Figure 36 shows a plot of arr/YS vs. CVN absorbed

energy for the three processing conditions examined in this report

and other data on as-rolled bainitic steels whose YS is similar to

the SAE 1141-Nb steel (52) . The CVN absorbed energy at room

temperature for the two types of steels, a-P and bainitic, are

similar, but the toughness in the tensile test is consistently

higher for the bainitic steels compared to that of the three

deformation conditions evaluated here for SAE 1141-Nb.

The equipment described in this report should produce results that

are compatible with the existing, small-scale TMP simulator at

NIST. A direct comparison between the results from the two systems

can be made for the condition D3 (deformed at 1100°C) . The recrystallized prior austenite grain size d after a given true strain

e can be predicted from the equation shown below:

drex = °* 9 ' do

°' 67 ' e

'°' 67

/ ( 15 ) where do is the starting austenite grain size prior to deformation.

Predictions from Eq. 15 agree well with the results of the small

simulator (46)

The measured value for condition D3 from this

study, 20 /im, agrees well with that predicted from Eq. 15, 16 /xm.

Fig. 35. Plot of strength in unaxial tension vs. hardness for a-P

steels tested in this study and ref. 36.

50
Fig. 36. Plot of tensile test result, ratio between the

interfacial stress at fracture arr and YS ay, and the energy

absorbed during CVN testing at room temperature. Data are from

this study and ref. 52.

3 . 5 Summary

1. The new, large-scale forging simulator at NIST is capable of

producing mechanical test data for the evaluation of the

structure-property relationships in steels.

2. Results from the new, large-scale simulator are consistent

with those produced on the existing small-scale TMP

simulator.

3 . Existing correlations between microstructure and properties

for a-P steels can predict YS, UTS, and % RA, but the

difference between measured and predicted can be as much as

100 MPa.

4. The evaluation of toughness is a major issue in the application of a-P steels. Impact testing is not a reliable indication of performance in many important industrial applications

where a-P steels could be used. A major component of our

research here at NIST should be directed at the development

51
of measurement technology for the evaluation of toughness

which is consistent with the application ( f itness-for-purpose) and future modeling efforts.

4. ACKNOWLEDGEMENTS

The material used in this study was supplied by the Chrysler

Corporation. We also want to thank James Hoffmann and Raymond

Filar of Chrysler Corporation, John Ham and Prakash Babu of

Caterpillar Tractor, William Heitmann of Inland Steel, and Michael

Burnett of Timken Corporation for their help and advice in the

preparation of this document.

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56
48. J.P. Houin, A. Simon, and G. Beck, "Relationship between

Structure and Mechanical Properties of Pearl ite between 0.2%

and 0 . 8%C, " Transactions ISIJ, Vol. 21, Laboratoire de

Metallurgie Associe au CNRS, Cedex, France, 1981, pp. 726-731.

49. N. Ridley, M. A. Malik, and G.W. Lorimer, "Partitioning and

Pearlite Growth Kinetics in an Ni-Cr Eutectoid Steel,"

Materials Characterization, Vol. 25, Elsevier Science

Publishing Co., Inc., New York, 1990, pp. 125-141.

50. R.M. Kuziak and Y. -W. Cheng, "The High-Temperature, High-Strain

Rate Flow Behavior of a Nb-Treated Bar Steel," to be published

in proceedings of International Symposium on Processing,

Microstructure, and Properties of Microalloyed, and Other

Modern High Strength Low Alloy Steels, Pittsburgh, PA, TMS

,

June 4-6, 1991.

51. Y.-W. Cheng, "Continuous-Cooling-Transformation Diagrams of a

Nb-Treated SAE 1141 Steel," to be published in proceedings of

International Symposium on Processing, Microstructure, and

Properties of Microalloyed, and Other Modern High Strength Low

Alloy Steels, Pittsburgh, PA, TMS, June 4-6, 1991.

52. D.A. Shepherd, work in progress, M.S. Thesis, Colorado School

of Mines, Golden, CO, 1991.

57
APPENDIX A

Stress Analysis

Argon et al. (Al) found that, for extremes in strain-hardening

behavior of a model material, the interfacial stress (cr rr ) at an

isolated particle could be approximated as

°rr = °o + aT < (!)

where aT = the hydrostatic stress at the inclusion and o

0 = the true

or equivalent plastic flow stress. The presence of the inclusion

raises the local stresses across the interface between particle and

matrix over the hydrostatic stress by an amount that approximately

equals the average plastic flow stress in the region. Their

studies suggest that when orr reaches a critical value ac

, void

nucleation occurs.

The approach described above can be applied at any point on the

stress-strain curve. Prior to necking, oQ

, is equal to o2

, the

measured load divided by the actual cross sectional area. After

maximum load is reached, the value of o

Q

modifies a

2

to account for

the change in stress state in the neck of the tensile specimen as

follows ( A2

)

o z [1+ (2R/ a) ] [In (l+a/2i?) ]

7

where R = the radius of curvature of the neck along the tensile

axis of the specimen, and a = the radius of the minimum cross

section of the specimen.

The value of <j t that develops at the particle is a function of o

0

and reflects the stress concentration due to necking. The ratio

Oj/o0 can be described by (A3) — = l/3+ln(l+a/2i?-r2 /2ai?) . (3)

o„

The value of r is the distance from the center of the diameter to

the specific point of interest in the cross section. The term

r

2/2aR in Eq. (3) was introduced to take into account the stress

distribution through the cross section. In this case, all

calculations were made in the center of the cross section where

r = 0 and the hydrostatic stress is a maximum.

58
References

Al. A.S. Argon and J. Im, Metallurgical Transactions A,- Vol.

6A, 1975, p. 839.

A2 . G. Dieter, in Mechanical Metallurgy . McGraw-Hill, New York,

1976.

A3. A.S. Argon, J. Im, and A. Needleman, Metallurgical

Transactions A, Vol. 6A, 1975, p. 815.

59
Mil

H