Heat Treatment and Toughness Behavior of Tool Steels (D2 and H13) for Cutting Blades

Attaullah. (Ayooq) Arain

Theses submitted in conformity with the requirernents for the degree of Master of Applied Science

Graduate Department of Metallurgy and Material Science University of Toronto

O Copyright by A. (Ayooq ) Arain 1999

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Acknowledgements

The author would like to express his sincere gratitude to his supervisor, Professor

Zhirui Wang, for his advice and encouragement throughout the course of this thesis.

Special thanks are due to Mohamrnad Hasnat (President) and George Kodama

(Technical Director) of A& M Heat Treat Ltd. for their assistance with the use of the

Vacuum Furnace and laboratory facilities, and for thought provoking discussions

pertaining to the data generated by this study. Material supplied by the Central

Welding Ltd is highly appreciated. Thanks are also due to Mr. F. Nueb and Sal

Boccia for their assistance in operating the SEM, and the author's research group

members, Dr. Bo Gong, Dr. Yang, Dr. Hamid S, Mr. Hai Ni, and Mr. John Yan, for

valuable discussions and collaborations.

Heat Treatment and Toughness Behavior of

Tool Steels (D2 and H13) for Cutting blades

Attaullah. (Ayooq) Arain

Department of Metallurgy and Materials Science

University of Toronto

ABSTRACT

The effects of austenitizing and tempering temperatures on the microstructure,

as-quenched and tempered hardness capability, and Charpy V-notch impact

resistance of D2 and H l 3 tool steels were investigated. Decarborization behavior

of D2 and Hi3 tool steels was observed by heat treating the samples in vacuum

and normal furnaces. Heat treatment in an open atmosphere furnace gave up to

.018"(.45rnm) thick layer of decarborization and also results in the loss of

precious alloying elements, which should be in controlled amounts for 02 and

H l 3 tool steels. The main results can be summarized as: (1) An increase in

austenitizing temperature resulted in coarsening of the grain structure, increased

dissolution of carbides, increased asquenched and tempered hardness

capability, and decreased impact toughness; (2) Tempering three times in

comparison with two times after hardening in a controlled atmosphere furnace

gives an increase in Charpy impact toughness of up to 25%; (3) 02 and Hl3

steels hardened at 1038'~ followed by three temperings show relatively higher

Charpy impact values versus those treated with one or two temperings. The

failure mechanism of the impact tested 02 and Hl3 steels after heat treatment at

1 OZS*C, 1038'~, and 4 065 '~ followed by the tempering up to three times at

temperatures 205'~, 538'~, 593%, and 6 2 0 ' ~ was studied through using

Scanning Electron Microscopy. The resultant microstructure of D2 and H i 3

steels after the three tempering process gives better plasticity than after two

tem perings.

iv Table of Contents Table of Contents

Table of Contents

List of Figures vi

List of Tables x

3 . Introduction 1

Background

Tool Steels

Category of Tool Steels

Cutting Performance of Tool Steel

Chernical Composition of Tool Steel

Heat Treatment of Tool Steel

Microstructure of Tod Steel

High Carbon-High Chromium Cold Work Tool Steels

Hot Work Tool Steels

Heat Treatment of Hot and Cold Work Tool Steels

02 Cold Work Tool Steel

H l 3 Hot Work Tool Steel

Objective

2. Experimental

2.1 . Materiais

2.2. Heat Treatment and hardness measurements

Table of Contents v

Table of Contents

2.3. Heat Treatrnent in vacuum and control atmosphere furnaces 23

2.4. Sample preparation for optical microscopy 24

2.5. Charpy V-notch impact testing 25

2.6. Scanning Electron Microscopy 27

3. Results and Discussion

3.1. Heat Treatrnent in Vacuum and Control Atmosphere Furnaces

3.2. Heat Treatment in open atmosphere furnace

3.3. Heat treatment by Current Operation

3.4. Charpy V-Notch Dynamic Impact Test

3.5. Surface analysis by using Scanning Electron Microscope

3.6. Hardness measurement in HRC

3.7. Summary of Major Results

4. Conclusions

References

vi List of Figures - - List of Figures

Fig. 1

Fig. 2

Fig. 3

Fig. 4

Fig. 5

Fig. 6

Fig. 7

Fig. 8

Fig. 9

Fig. 10

Fig. 11

Fig. 12

List of Figures

Schematic of tool steel heat treatments.

Schematic continuous cooling diagrarn for a typical tool steel.

Isothemal section of the iron-chromium-carbon system at 700%.

Isothemal section of the iron-chromium-carbon system at 1 0 0 0 ~ ~ .

Vertical section for 5 pct Cr alloys of the Fe-Cr-C system.

Charpy V-Notch impact testing machine.

Optical micrograph of D2 tool steel heat treated in vacuum furnace at

1 02S°C/30min and tempered twice at 538OCI2hrs.

Optical micrograph of D2 tool steel heat treated in vacuum furnace at

1038"C/3Omin and tempered twice at 538"C/2hrs.

Optical micrograph of 02 tool steel heat treated in vacuum furnace at

1 065OC130min and tempered hnrice at 538"CIZhrs.

Optical micrograph of H l3 tool steel heat treated in vacuum furnace at

1025'C/30min and tempered Wice at 538'C and 593"C/2hrs.

Optical micrograph of Hl3 tool steel heat treated in vacuum furnace at

1065'C/30min and tempered Nice at 538% t2hn.

Microhardness measurement of 02 tool steel in Knoop scale by using

weight 500 gram.

vii List of Figures - - List of Figures

Fig. 13 Microhardness measurement of H l3 tool steel in Knoop scale by using

weight 500 gram.

Fig. 14 Optical micrograph of D2 tool steel heat treated by current operation

(open atmosphere) at 1025"C/30min and tempered twice.

Fig. 15 Optical micrograph of H l3 tool steel heat treated by current operation

(open atmosphere) at 1025"C/30min and tempered twice.

Fig. 16 V-Notch Charpy impact test result for H l 3 tool steel heat treated at

1025"C, 1038OC, 1065OC/30min. and tempered up to three times at

538°C and 593OC12hrs.

Fig. 17 V-Notch Charpy impact test result for Hl3 tool steel heat treated at

1025"C, 1038OC, 1065OCf30rnin. and tempered up to three times at

538°C 12hrs.

Fig. 18 V-Notch Charpy impact test result for H13 tool steel heat treated at

1038*C, 1065°C130min. and tempered up to three times at 538°C and

62O0C12hrs.

Fig. 19 V-Notch Charpy impact test result for D2 tool steel heat treated at

1025OC, 1038OC, 1065OC/30min. and ternpered up to three times at

593OC12hrs.

Fig. 20 V-Notch Charpy impact test result for D2 tool steel heat treated at

1025OC, 1038OC, 1065°C130min. and tempered two times at 20S°C/2hrs.

Fig. 21 V-Notch Charpy impact test result for D2 tool steel heat treated at

1025*C, 1038OC, 1065°C/30rnin. and tempered two times at 538OCMhrs.

viii List of Figures - - List of Figures

Fig. 22

Fig. 23

Fig. 24

Fig. 25

Fig. 26

Fig. 27

Fig. 28

Fig. 29

SEM micrograph of D2 tool steel heat treated in controlled atmosphere

fumace at 1025'C/30min and tempered at 593'C 12hrs.

SEM micrograph of D2 tool steel heat treated in controlled atmosphere

furnace at 1038"C/30min and tempered at 593°C 12hn showing the

matching part and morphology of fractured surface.

SEM micrograph of 02 tool steel heat treated in controlled atmosphere

furnace at 1038"C130min and tempered at 593°C 12hrs showing the

matching part of fractured carbide.

SEM rnicrograph of 02 tool steel heat treated in controlled atmosphere

furnace at 1065'C130min and tempered thrice at 593% nhrs showing

the morphology of fracture.

SEM micrograph of H l3 tool steel heat treated in controlled atmosphere

furnace at 1025"C/30min and tempered thrice at 538OC, 593°C and

593°C /2hrs showing the difference between two and three tempering.

SEM micrograph of Hl3 tool steel heat treated in controlled atmosphere

fumace at 1065'C/30min and tempered thrice at 538OC, 593°C and

593°C 12hrs showing the morphology of fracture surface.

SEM rnicrograph of Hl3 tool steel heat treated in controlled atmosphere

fumace at 1025"C/30min and tempered at 538°C and 593°C /2hrs

showing the difference between two and three tempers.

Hardness data (HRC) for V-Notch Charpy impact 02 tool steel samples

heat treated at 1025OC, 1038OC, 1065OC130min. and tempered up to

three times at 593OCt2hrs.

ix List of Figures - List of Figures

Fig. 30

Fig. 31

Fig. 32

Fig. 33

Fig. 34

Hardness data (HRC) for V-Notch Charpy impact 02 tool steel samples

heat treated at 1025OC, 1038OC, 1065°C130min. and tempered up to

three times at 538°C/2hrs.

Hardness data (HRC) for V-Notch Charpy impact D2 tool steel samples

heat treated at 1025°C1 1038OC, 1065°C/30min. and tempered up to

three tirnes at 205°Ci2hrs.

Hardness data (HRC) for V-Notch Charpy impact H l 3 tool steel samples

heat treated at 1025OC, 1038OC, 1065OC/30min. and tempered up to

three tirnes at 538OC and 593OCI2hrs.

Hardness data (HRC) for V-Notch Charpy impact H l 3 tool steel samples

heat treated at 1025OC, 10380C1 1065°C/30min. and tempered up to

three times at 538OC/2hrs.

Hardness data (HRC) for V-Notch Charpy impact Hl3 tool steel samples

heat treated at 1025OC, 1038OC, 1065OC/30min. and tempered up to

three times at 538°C and 620°C/2hrs.

List of Tables List of Tables

List of Tables

Table 2.1. Showing austenitizing temperature, soaking time, and cooling pradice

used for 02 and H 1 3 tool steel.

Table 3.1. HRC and V-Notch Charpy impact values for 02 steel at several

austenitizingand tempering temperatures.

Table 3.2. HRC and V-Notch Charpy impact values for H l3 steel by using

several austenitizing and tempering temperatures.

1 Introduction

1.1 Background:

Research work was carried out on 02 and Hl3 tool steels to increase the

service life of blades and knives manufactured by the Central Welding Ltd, which

is the only Canadian manufacturer of heavy-âuty shear blades and machine

knives, used in the steel and scrap metal industries. This Company is specialized

in the manufacturing of rotary edge trimmers and slitter knives. At the present

time, premature replacement of D2 and H 1 3 tool steel knives is often required as

the result of cutting edge chipping and dulling. Chipping is a phenornenon in

which srnall parts of the material break away from the cutting edge. Chipping

apparently will damage the integrity of the cutting edge counter leading to a poor

cutting process and unsatisfactory products. Dulling is the phenornenon in which

the sharp edge of the knives or blades becorne "rounded" after a certain period of

normal use, Le., application without overloading or over-pressure. In order to

increase the service life of the blades and knives, research on 02 and Hf3 tool

steels has therefore been conducted supported by OCMR and Central Welding

Ltd .

1.2 Tool Steels

A tool steel is any steel used to make tools for cutting, foming or otheiwise

shaping a material into a part or component adapted to a definite use. The

addition of relatively large amounts of tungsten, molybdenum, manganese and

chromium can enable tool steels to meet stn'ngent service demands and can

provide greater dimensional control and freedom from cracking during heat

treatment.

The performance of a tool in senrice depends on the design of the tool, the

accuracy with which the tool is made, the choice of tool steel, and the choice of

heat treatment. High quality tool steel, appropriate design, and proper

manufacturing methods are the essential factors determining the procedure of

the heat treatment.

1.2.1 Categoiies of Tool Steels

It is important to classify tool steels into a relatively small number of groups for

purposes of comparison and evaluation and to facilitate the selection of steel for

a particular application. Because tool steels are of such diverse compositions, it

has never been easy to fit them into one category of the alloy steel system. Tool

steels have narrow Iimits on the amounts of alloying elements, and entire series

of steels are based on the variation in carbon content. The methods used most

frequently for classification of tool steels are the "Society for Automotive

Engineers" (SAE) and "American lron and Steel Institute" (AISI) rnethods. The

AISI method is more popular because it makes tool steel classification more

simple and understandable.

The AISI classification of tool steels will be used throughout the balance of this

study. The classification and a brief sumrnary of the major features of each class

is mentioned below (1,2):

Water-hardening tool steek, type - W

Shock-resisting tool steels, type - S

Mold tool steels, type - P

Special-purpose tool steels, type - L and F

Cold work tool steels, type - O, A and D

Hot work tool steels, type - H

High-speed tool steels, type - T and M

The water hardening tool steels, AISI type W, have the lowest alloy content and

therefore the lowest hardenability of any of the tool steels. As a result, the W tool

steels frequently require water quenching and heavy sections harden only to

shallow depths. Thin sections can be hardened by oil quenching to minimize

quench cracking and distortion.

The shock-resistant tool steels, AISI type S, have lower carbon content and

somewhat higher alloy content than the W steels. The medium carbon content

improves toughness and makes the type S steels good for applications with

shock and impact loading.

Tool steels for cold work include three classes of steels, AlSI type O, A and D.

These classes each have high carbon content for high hardness and high Wear

resistance in cold work applications, but differ in alloy content, which affects

hardenability and the carbide distributions incorporated into the hardened

microstructures.

Tool steels used for dies to mold plastics, AlSI type P. are exposed to less

severe Wear than metal-working steel, and therefore have low carbon content. A

key requirement is good polishability and excellent surface finish.

Hot work tool steels, AISI type Hl fall into groups which have either chromium,

tungsten, or molybdenurn as the major alloying element. The H steels are used

for hot forging, metal shearing, and metal die-casting dies.

The high-speed tool steels are highly alloyed, with tungsten, and molybdenum

as the major alloying elements in the T and M grades, respectively. The tungsten

and vanadium in these steels produce very high densities of stable carbides. As

a result, the high-speed tool steels are capable of retaining hardness at high

temperatures and are widely used for high-speed cutting and machining

applications.

1.3 Cutting performance of tool steel

Normally. cutting performance refers to tool life until the tool is reground or until

end-wear. Cutting performance is econornically significant because production

costs are influenced by it. Cutting performance of tool steel can be judged by

mechanical properties such as sharpness, hardness, strength, toughness, and

microstructure of tool steel. These mechanical properties, however, are

influenced by the chernical composition and heat treatment of the tool, which

affect the tool's microstructure.

1.3.1 Chernical composition of tool steel

Chernical composition is the most important influence upon shearing

performance of the tool steel. Each alloying element in tool steel. such as

tungsten, chromium, molybdenum, vanadium, has a specific role in detemining

5

the mechanical properties. For hot work tool steels and high-carbon high-

chromiurn tool steels, slow cooling du ring solidification results in large amounts of

segregate that is a carbide of different alloys and is deposited frorn the melt as a

eutectic mixture of austenite and carbide. After solidification, such a segregate

can be broken up only with difficulty and then, only by mechanical work. These

carôides are brittle and their nonunifonn distribution causes the steel to possess

limited ductility and also variation in chemical composition. Therefore, it becomes

highly important in the "freezing" of the steel that the distribution of the carbide

segregate be as closely controlled as possible. Also mechan ical working

operations should be strictly controlled to avoid change in chemical composition.

1.3.2 Heat treatrnent of Tool Steel

The heat treatment to which a tool has been subjected has a marked influence

on cutting performance of tool steel (3). The general heat treatment schedules

applied to tool steels are shown in figure 1. Austenitizing is a very critical step in

the hardening of tool steel. It is in this step that the final alloy elements are

partitioned between the austenitic matrix (which will transfonn to martensite) and

the retained carbides. This partitioning fixes the chemistry, volume fraction, and

dispersion of the retained carbides. The retained alloy carbides not only

contribute to Wear resistance, but also wntrol austenitic grain size. The finer the

carbides and the larger the volume fraction of carbides, the more effectively

austenitic grain growth is controlled. If during heating the austenitizing

temperature is high, the carbide will dissolve to a large extent, and the

Figure: 1

TIME

Schematic of tool steel heat treatment (4).

Figure: 2 Schematic continuous cooling diagram for a typical tool steel. Tl, T2, and T3 npresent decieasing cooling rates, Ci, Pi and Bi repmsents the initiation of carbide, pearlite, and bainite formation respectively (5).

-

precipitation of cementite on cooling will have a greater tendency to take place at

coarse austenite grain boundaries. If, however, the carbide has not been

completely dissolved and brge quantities remain in the fom of rounded particles

throughout the matrix, carbide precipitation will take place on these preexisting

points, and the network of cementite surrounding the grain boundary will not

form. Thus, overly high austenitizing temperatures rnust be avoided so as to

prevent grain growth which can led to problems with cracking, retained austenite,

and excessive distortion.

Relatively slow oil quenching or air cooling for hardening of tool steels can lead

to grain boundary carbide formation, which makes tool steel susceptible to

intergranular failure. Figure 2 shows schematically the effects of three cooling

rates on the transformation of a typical tool steel (5). The high hardenability of

tool steels effectively suppresses perlite formation at al1 cooling rates. Bainite

formation is also readily suppressed except in heavy sections, which cool slowly.

However by slow cooling, the formation of carbides on austenite grain

boundaries is difficult to suppress, as shown in fig 2. Small amounts of carbides

do not significantly affect hardness but may lower tool steel fracture resistance,

lead ing to quench cracking, intergranular fracture of tool steels and reduced

performance of hot work tool steels such as Hl3 (6,7). A number of

investigations have shown that the presence of the grain boundary carbides

significantly reduces toug hness of hardened and tempered tool steels (6,8,9).

Tool steels hardened in an oxidizing atmosphere scale freely; the 02 and Hl3

tool steels cannot be hardened in this manner without excessive decarborization .

Steels that can be hardened satisfactorily in an oxidizing atmosphere generally

have low chromium content ( 1% or less ) and do not require a high hardening

temperature ( no greater than 870°C ) (44). In order to protect the tool steel

surface from decarborization and scaling during heat treatment, the fumace

medium must be kept neutral. Othewise decarborization will lead to soft surfaces

and cause cracking due to the formation of residual tensile stresses in the

surface. A possible explanation of this mechanism is that the reductions in

carbon content raise the rnartensite transformation temperature. Thus, on

quenching, the outer layers transfomi first at a much higher temperature, and

when the core transforms and expands, it puts the outer layer in tension (1 0).

1.3.3 Microstructure of Tool Steel

The cutting performance of heat treated tool steel can be improved by obtaining

finer grain size, a minimum amount of retained austenite, spheroid and finer

carbide size and a unifom distribution of carbides (3). As mentioned above, the

austenitizing temperature and quenching time should be appropriate, otherwise

grain growth and an increased amount of retained austenite and segregation of

carbides along grain boundaries will occur and can significantly reduce the life

and cutting performance of tool steel.

1.4 Hig hcarbon hig h-chromium cold work tool steels

In general, high-carbon high-chromium steels can be divided into Mo main

categories: those that are essentially oil hardening and those that are essentially

air hardening. Further division can also be made on the basis of carbon content.

The original high-carbon high-chromiurn steels contained frorn 2.00 to 2.50%

carbon. Later modifications to obtain better machinability and less brittleness

lowered the carbon content to the range of 1.00 to 1.50% (2). Cold-work steels

should show the following physical characteristics:

1. Low rnovement in hardening .

2. High asquench hardness.

3. Maintenance of a keen edge for cutting purposes.

4. Resistance to mechanical shock.

5. Good machinability in annealed condition.

The nominal composition of high-Carbon highChromium cold work cutting

steels is: C: 1.5 to 2.00%, Mn: .30 Oh, Si: 0.25 to 0.85%, Cr: 12%. V: 0.25 to

0.6%, and Mo: 0.5 to 1%. The outstanding characteristics of cold work cutting

steels are high hardenability, Wear resistance, and high strength.

lsothermal sections of the ternary iron-chromium-carbon system provide insight

into the structure and properties of chromium cold work steels. Figure 3 shows

an isothemal section through this system at 700°c, which should correspond

closely to the room temperature section. Therefore, this diagram represents

phases present in annealed alloys. When the ratio of chromium to carbon

exceeds 3:1, chrornium-rich carbides are found ((CrFe)& or (CrFe)7Ca or

both). In high-carbon high-chromium steels, (CrFe)7C3 carbide is predominant. In

more complex steels, the metal lattice of the carbides may also contain

molybdenum, vanadium, silicon, and manganese (45,46). The composlion of

Figure: 3 lsotheimal section of the ironthromium-carbon system at 700 '~ (1 1).

- -p -

carbides found in this system is variable and depends on the overall composition

of the alloy. Molybdenum or tungsten present in some of these steels stabilizes

the (CrFe)23C6 carbide. For example, D2 with molybdenum higher than usual

(1.41 % Cl 13.1 3% Cr, 1.2% Mo) is reported to contain only (CrFe)& carbide in

the annealed condition (12). Most of the vanadium, some of the cobalt, but

relatively little of the nickel that may be present are also found in the carbide

phase (13). Each carbide present in steel has a different nature, and Wear

resistance of the steel is detenined by the amount and nature of that carbide:

the harder the excess carbide, the greater the Wear resistance. Microhardness

measurements made using the Knoop scale found the (CrFe)7C3 carbide to be

considerably harder (1820 Knoop) than the cementite in a plain carbon tool steel

(1 150 Knoop) (1 4, 15).

On austenitizing some of the carbide dissolves in the austenite, thus supplying

the matrix with alloying elements necessary for high hardenability and as-

quenched hardness. The isothermal section shown in Figure 4 dernonstrates the

structures prevailing at the austenitizing temperature (1 0 0 0 ~ ~ ) . (CrFe)7C3 carbide

is the only carbide present in high-carbon highthromiurn steels after

austenitizing at 1 000' C.

The presence of high chromium content enables these steels to resist oxidation

at high temperatures to a much greater degree than carbon or other low-alloy

steels. High chromium content also causes an appreciable resistance to staining

when the steel is hardened and polished.

Figure: 4 lsothermal section of the iron-chromium-carbon

Other elements added in small amounts to high-carbon high-chromium steels

include vanadium, cobalt, silicon, and tungsten. Yamanaka (17) has studied the

effect of variations of molybdenum (O to 1.5%) and vanadium (O to 1.2%) on the

properties of D2 tool steel and found that molybdenum increases hardenability

and toughness but has little effect on austenite grain size or the quantity of

retained austenite. Vanadium (in proportions greater than 0.8%) produces fine

grain size but decreases hardenability (the austenïtizing temperature required to

produce full hardness increases wlh increasing vanadium). Vanadium decreases

retained austenite and, with proportions up to 1 %, improves toughness.

Typical applications of high-carbon high-chromium cold work tool steels include

shear blades, slitting cutters, cold extrusion dies, punches, broaches, mandrels,

forming and bending rolls, and hot trimming of forgings.

1.5 Hot work tool steels

In general, hot work steels are of the medium and high-alloy type, and most of

them have relatively low carbon content (0.25 to 0.6%). Hot work steels should

show the following physical characteristics:

1. Resistance to deformation at the working temperature.

2. Resistance to shock.

3. Resistance to Wear at the working temperature.

4. Resistance to heat treating deformation.

5. Resistance to heat checking.

6. Good machinability in the annealed condition.

The nominal composition of chromium-molybdenum hot work cutting steels is:

C: .35 to .40%. Mn: .30 to .60%, Si: 1.0%. Cr: 3.50 to 5.0%, V: .4 to 1.0%,

W: 1.25 to 1.50%, and Mo: 1.50 to 2.5%. The outstanding characteristics of hot

work cutting steels are toug hness, shock resistance, and hot hardness.

Typical applications of chromium-molybdenum hot work steels include die-

casting dies. forging dies, shear blades for hot work, punches, piercers and

mandrels for hot work. hot extrusion tooling, and al1 types of dies for hot work that

involves shock. Certain of these steels are used for ultra high-strength structural

parts (18).

Chromium-molybdenum steels have extremely high hardenability. Molybdenum,

which is present in an amount of 1% or greater, is responsible in large rneasure

for this property. Tungsten, which may be present, contributes little to

hardenability, and vanadium actually decreases it by tying up carbon in the form

of stable vanadium carbides. The high silicon content in these steels improves

oxidation resistance while changing the type of scale formed on air cooling to

enable its easy removal. Either carborization or decarborization of these steels

increases the tendency to heat checking. The vertical section (Fig. 5) (shows

temperature ranges over which the various carbides coexist with austenite and

ferrite) for Fe-Cr-C alloys containing 5 wt.% chrornium. This information is useful

in designing hot work schedules and heat treatments for annealing and

hardening (4,9,16).

Figure: 5 Vertical section for 5 pct Cr alloys of the Fe-Cr-C system. Vertical dashed lines indicate phase equilibria of the alloys based only on their chromium carbon contents. A, F and L designate austenite, ferrite and liquid respectively (39).

1.6 Heat treatment of cold and hot work tool steels

Heat treatment of tool steels for cutting purposes are conducted to produce an

optimal combination of high hardness, good Wear resistance, and sufficient

fracture resistance or toughness for a given application. High hardness is

frequently produced by the transformation of austenite into martensite, and

toughness is controlled largely by the tempering of the martensite.

Heat treatment to produce martensite consists of three steps: heating to the

austenitizing temperature, austenitizing, and cooling or quenching. Heating of the

ferrite-sp heroidized carbide microstructure to the austenitizing temperature in

highly alloyed tool steels requires a preheat step to equalize the ternperature

through a section, thereby preventing distortion or cracking that might occur if the

surface and centre sections heat at significantly different rates (19,20).

Once the austenite is fomed, the alloying elements and carbon partition

themselves between the austenite and the carbides according to the

requirements of equilibrium at a given temperature. As the carbides dissolve, the

austenite becomes rich in carbon and alloying element content (21-23). The

austenitizing of cold and hot work cutting steel is designed to retain a significant

volume fraction of spheroidized carbides for the following purposes: to produce

austenite of optimum composition; to improve Wear resistance during service;

and to prevent grain coanening and abnomial grain growth during austenitizing

(40)-

The hardenability of cold and hot work cutting tool steels is quite high, and

therefore the steels can generally be hardened by air cooling. When diffusion

controlled transformations are avoided, the austenite remains untransfotmed until

the martensite start temperature (M,) is reached. At this temperature, the

diffusionless transformation of austenite into martensite begins. The higher the

carbon and alloy content of the austenite, the lower the M, temperature (24),

which results in less martensite formation at room temperature. Thus in highly

alloyed austenite, a considerable fraction of the austenite might be retained in the

microstructure at room temperature, resulting in a much lower hardness than

expected for more com pletely transformed microstructures. When retained

austenite is of concern, the austenitizing temperature can be decreased to retain

more carbides. Subzero cooling to transform additional austenite into martensite

is also used sometimes (20). If sewndary hardening is required during

tempering, austenitizing is designed to put as much alloy and carbon in solution

as possible, while avoiding abnormal grain growth and excessive retained

austenite.

1.7 02 tool steel

D2 tool steel used for cutting purposes operates under conditions of impact,

where resistance to mechanical damage is desired. Due to high carbon and high

chromium content, the Wear resistance of D2 tool steel is approxirnately eight

times that of plain carbon steels (25). The chernical composition of 02 steel is

usually: C-1.5%, Mn-0.30%, Si-0.25%, Cr4 2%, V-0.60%, Mo-0.80%.

Kligler (26) has shown that the mechanical properties of D2 steel are

anisotropic and depend on orientation with respect to the rolling direction. 60th

strength and ductility, as measured by tension, compression and bend tests were

found to be maximum in the direction parallel to the rolling direction and minimum

in the direction transverse to the rolling direction. This directionality of mechanical

properties can be attributed to the production of eiongated carbide stringers in

the direction of rolling.

The dimensional changes resulting from hardening of high-carbon high-

chromiurn steels are exceptionally small. Previous research (27-29) showed that

an 1 1.00% chromium steel of this type expanded only 0.1 % of the annealed

volume after hardening in air.

Although the majority of applications of D2 tool steel involve cold work, it is also

widely used for hot trimming of forgings. Typical applications include blanking

dies, slitting cutters, shear blades, forming dies, knurls, gages (plug and thread),

punches, trimming dies, etc.

1.8 Hl3 tool steel

Hi3 tool steel that belongs to the hot working tool steel group is the most

frequently used steel in this group. This steel possesses a combination of hot

strength, Wear resistance and toughness, and is predominantly based on the

0.4%C, 5%Cr compositions containing up to l.S%Mo, 1 %V and sometimes with

increased silicon.

To maintain the required properties at high temperatures in Hl3 tool steel, the

most convenient method is to use a secondary hardening reaction involving the

precipitation of alloy carbides such as Mo2C and VC after tempering at

5 0 0 ~ 1 6 5 0 ~ ~ .

The most stable carbide in H l3 steel is VC, which mostly remains undissolved

at recommended austenitizing temperatures. These undissolved, uniformly

distributed carbide particles are pinned to the austenite grain boundaries and

help to maintain the fine austenite grain size. H l 3 is an air hardening steel, but at

the slower cooling rates in larger sections there are increased amounts of both

lower and upper bainite (30), and also an increased tendency for carbides to be

precipitated du ring cooling on the austenitite grain boundaries. It is well known

(30-32) that upper bainite impairs both the high ductility and impact toughness,

and a similar detrimental effect is also produced by grain boundary carbides

which are mainly of the VC type (32).

During tempering of Hl3 tool steel, secondary hardening occurs due to

precipitating carbide being VC in which some molybdenum is dissolved (23).

Because secondary hardening is due to precipitation, its intensity increases with

increasing volume fraction and decreasing particle size of the alloy carbide.

Austenitizing at hig her temperature provides a greater num ber of nuclei du ring

tempering, and consequently a smaller particle size, smaller interparticle spacing

and greater intensity of secondary hardening. It has been suggested (34.35) that

at lower austenitizing temperatures there are VC clusters which are inherited by

the martensite and act as nuclei for VC precipitation during tempering. Hence

coarser precipitates are formed, and if the VC clusten segregate to austenite

grain boundaries, more VC would fonn on the boundaries during tempering and

-

give even less secondary hardening. On the other hand, austenitizing at higher

temperatures causes the VC clusten to be thermally dispersed so that the VC

precipitated during tempering has no precipitation nuclei and forms a greater

number of smaller particles with a consequent greater intensity of secondary

hardening and a higher overaged hardness. Thus, increasing the austenitizing

temperature not only dissolves more VC and gives a larger volume fraction of

precipitate during tempering, but also refines the precipitate particle size. Both

effects lead to greater secondary hardening and higher overaged hardness. Of

course, the austenitizing temperature must not be increased so much that grain

coarsening takes place. During overaging at high tempering temperatures, the

VC coarsens slowly and MZ3 C6 is precipitated, possibly at the expense of some

dissolution of VC (33,36). Hl3 tool steel is not tempered at maximum secondary

hardness due to a marked embrittlement which occurs (37-39).

1.9 Objectives

This project has following objectives:

1. To investigate failure mechanisms of the 02 and Hl3 cutting tools with

consideration of the relationship between current heat treatment procedures

and the resultant microstructures/hardness in these materials.

2. To evaluate the dynamic impact toughness at several heat treatment

conditions by using Charpy V-notch impact toughness tests for 02 and Hl3

tool steels.

3. f o study the fracture surface of the Charpy impact tested specimens using

Scanning Electron Microscopy (SEM) in order to understand the failure

mechanism.

4. To determine the optimum conditions for heat treatments of these materials

for which the properties and life of the tools can be improved. Parameters

considered include austenitizing temperature, tempering temperature. and

number of temperings.

2. Experimental

2.1. Materials

The materials used in this investigation were supplied by Central Welding Ltd. in

annealed conditions. Chernical analyses were performed on the 02 and Hl3

steel samples used in this study. The samples were: commercial grade tool steel

D2 with a nominal composition of 1.49 wt. % carbon, 11.59 wt. % chromium, .80

wt. % molybdenum. .80 wt. % vanadium, .43 wt. % manganese. .39 wt. % silicon.

and .14 wt % nickel; and commercial grade tool steel H l3 with a nominal

composition of .38 wt % carbon, 4.89 wt. % chromium, 1.32 wt. % molybdenum,

1.11 wt. % vanadium, .99 wt. % silicon, .35 wt. % rnanganese, and .36 wt %

nickel. The results from these analyses indicated that both D2 and Hl3 steels

were within the limits set forth by AlSI (51 ). Samples from prematurely failed D2

and H l 3 tool steel cutting blades were also used in this study.

2.2 Heat treatment and hardness measurements

In order to study the behaviour of D2 and Hl3 tool steels at various

austenitizing temperatures, annealed samples of 02 and H l 3 tool steels were

heat treated in a vacuum, a controlled atmosphere, and an open atmosphere

furnace. After hardening, several tempering temperatures were used to

investigate the relation between hardness, microstrudure, and impact toughness.

The samples (1 in x 1 in x 1 in) were cut from the rolled and annealed plates before

they were heat treated. After heat treatment, the heat treated samples were

ground on abrasive paper to 600 grit. Then hardness tests were carried out on a

Rockwell tester and a micro hardness tester using the C-scale (HRC) and

Knoop-scale (KHN), respectively. The standard penetration was obtained with a

120' sphero-conical diamond indentor on the Rockwell sa le and with a rhombic-

based pyramidal diamond on the Knoop scale, and the applied major loads were

150 kgf and 500 grams, respectively. The hardness tests provided a measure of

D2 and Hl3 tool steel resistance to permanent or plastic defornation after

hardening and tempering at various temperatures. At least four hardness

measurements were made for each heat treatment condition to ensure accurate

results.

2.3 Heat treatment in vacuum and controlled atmosphere

furnaces

D2 and Hl3 tool steel annealed samples were heated slowly and uniformly to

the austenitizing temperatures 1 02S°C, 1 038'C, and 1065°C in a horizontal box-

type vacuum furnace with heating on two sides and gas cooling from bottom to

top. The maximum nitrogen cooling gas pressure at quenching was 2 Bar. An

electrically heated furnace with argon protective atmosphere was also used to

austenitize the D2 and Hl3 tool steel samples at the same above mentioned

temperatures and followed by air cooling. After cooling, the D2 and H l 3 tool steel

samples were tempered in order to release stresses that developd during

quenching and also to obtain optimum toughness. 60th furnaces were calibrated

before and during heat treatment by using extra thermocouple.

2.4 Sample preparation for optical microscopy

Metallographic preparation of 02 and Hl3 tool steels is relatively difficult due

to the large amount of carbides. The test pieces were cut from the annealed and

heat treated samples using an abrasive cutting wheel. Due to the large amount of

massive carbide particles, it was very difFicult to cut the samples, even when in

annealed condition, so extreme care had been taken during the cutting process.

A soft grade of cutting wheel, a copious supply of coolant, and a slow cutting

speed were used to avoid over heating and breaking the carbide particles. Each

of these effects can lead to misinterpretation of the observed microstructure. A

Buehler mounting machine was used to mount the sarnples, which were attached

with Transoptic mounting powder; a pressure of up to 3000-psi and a

temperature of up to 6 6 ' ~ were used during mounting.

After rnounting, grinding and polishing of the specimen was carried out in

several steps. Motor-driven disk grinders were used with 240, 320,400: 600, and

1200 grit grinding papers. After fine grinding, polishing produced a surface that

was Rat, scratch-free, and rnirror-like in appearance. For mechanical polishing,

the specimen was introduced to a cloth, and 1pm alumina particle spray was

introduced for a short time at a low initial pressure. The pressure was increased

for the main polishing time and then reduced toward the end. To polish the

scratch-free surface, a .5pm diamond paste was applied in high concentration at

the beginning, and smaller quantities were applied as required during the

polishing stage. The polishing time was kept short and the pressure was kept

25

low. A long polishing time or high pressure can result in the formation of relief,

because of the carbide particles, or may pull out or drag inclusions. A

microscopie examination of the surface s hould not reveal any polishing scratches

or any residual abrasive. All the samples were observed under an MeF3 optical

microscope.

2.5 Charpy V-Notch Impact Testing

One hundred and fi@ Charpy impact test blanks (20 mm square x 70mm) were

saw cut from square bars of 02 and H l 3 tool steels, respectively. The specimen

blanks were machined to diameter of 10 mm by a finish length of 55 mm

(tolerance t .O5 mm) with a notch of radius 22 X0 @or to heat treating. Twenty

five machined V-Notch samples of each of 02 and Hl3 were austenitized in a

controlled atmosphere furnace by using one of the three heat treatment cycles

outlined in Table 2.1. Following austenitizing, specimens from each 02 and H l 3

tool steel were single, double and triple tempered at either 205, 538. 593, or

620% for 2 hours + 2 hours and + 2 hours (For accurate results three specimens

were used at each tempering temperature) in protective atmosphere in order to

avoid any decarborization at the notch root of samples. Charpy impact testing

was conducted at room temperature by using The Hounsfield Balanced Impact

Machine, which has 48 ft-lb maximum capacity and is accurate to t .l ft-lb. The

working procedure of the V-Notch Charpy impact testing machine is shown in

Figure 6 (52).

Figure: 6 Charpy V-Notch impact testing machine. -

~iagram shoiwing- impact hammer, W, dropping from height, h l ,impacthg at C and rising to final height, h2. The energy absorbed, h2 - hl, is recorded on dial D (52).

Table 2.1 Austenitizing temperature, soaking time, and cooling pnctice

used for 0 2 and Hl3 tool steels.

Austenitking Temperature["CI 1 Soaking Time 1 Cooling Practice

1024°C

1038'C

Upon completion of this testing, the Rockwell C hardness of select impact

samples was evaluated in order to measure the tempered hardness of each set

of test specimens. In addition, the tempering response of these specimens

coupled with the Charpy V-Notch impact test data were used to evaluate the

strengthltoughness combination that was achieved in the specimens based on

the austenitizing parameten investigated in this study.

1 065°C

2.6 Scanning Electron Mictoscopy

In order to more thoroughly understand the effects of austenitizing temperature

and tempering behavior on the impact properties of 02 and Hl3 tool steels, the

fracture surfaces of selected specimens were examined using a Hitachi S-570

combined with a Link EDX system Scanning Electron Microscope (SEM)

operating at an accelerating voltage of 20 kV. Representative images of the

obsewed features on these fracture surfaces were recorded.

30min

30min

Air Cool

Air Cool

30min Air Cool

3. Results and discussion

3.1 Heat treatment in vacuum and controlled atmosphere

furnaces

After heat treatment, the samples of D2 and Hl3 steeis were observed under

optical microscope. The results revealed that the heat treatment in the vacuum

furnace and the control ted atmosp here furnace, at appropriate austenitizing

temperatures, gave uniforrn microstructure, and no decarborization and scaling

was observed. An increase in austenitizing temperatures will affect the

microstructure, austenite grain size, carbide solutioning behavior, and other

properties such as Charpy V-notch impact toughness, hardness capability and

temper resistance of D2 and Hl 3 tool steels (4).

Figure 7 shows the optical micrograph for 02 tool steel that was heat treated in

the vacuum furnace at 1025 '~ and tempered at 538 '~ . The microstructure

reveals tempered martensite in which coane carbides (The carbides those do not

dissolve during austenitizing) dispersion coexists with fine carbides (The carbides

those precepitate during tempering) dispersion. The coarse carbide particles are

expected to be M7C3 (chromium carbides) (4). Coarse carbides act as barriers to

austenite grain growth and are responsible to a large degree for the high Wear

resistance. The shape and distribution of these carbide particles are believed ta

be responsible for the anisotropic mechanical properües. The micrograph also

shows the fine carbide particles that are precipitated after tempering.

Fig: 7

02 Tool Steel heat treated in vacuum furnace at 1025°C/30min and tempered Nice at 538OC12HRS. ( M t ) 320X and (Right) 800X.

Fig 8 shows the optical micrograph for D2 steel hardened and tempered at

1038 '~ and 538 '~ , respectively. The microstructure reveals a fine grain size, not

much dissolution of carbides at this austenitizing temperature, and coarser

carbides with angular shape and finer carbides with spheroid shape existing

throughout the martensitic matrix. This observation is also supported by ref (2).

Austenitizing at 1065~~130min and tempering at 538O~12hrs (Fig 9) resulted in a

significant amount of the coarser and finer carbides dissolved into the matrix but

even coarser carbides are present to stop rapid grain growth. Therefore, there is

no evidence of a large amount of grain growth at this temperature, and study of

the micrograph reveals the carbides are precipitated along grain boundaries.

D2 and H l3 steel blades should have an optimum combination of high hardness,

good Wear resistance and sufficient fracture resistance or toughness for a given

application. The austenitizing of 02 and Hl3 tool steels should be designed to

retain a significant volume fraction of spheroidized carbides to produce austenite

in balance composition. The retained carbides also contribute significantly to Wear

resistance d uring service and unifon distributions of carbides are necessary to

prevent grain coanening and abnomal grain growth during austenitizing (40).

The optical micrograph of H l3 tool steel specimen heat treated at 1025'~ in the

vacuum fumace followed by tempering (Fig 10) shows martensite with small pre-

existing austenite grain boundaries and spheroid carbide particles that are

distributed throughout the matrix. These martensitic structures were very uniforni

Fig: 8

D2 Tool Steel heat treated in vacuum furnace at 1038OCJ30min and

tempe- twice at 538OC12HRS. 625X.

Fig: 9

02 Tool Steel heat treated in vacuum furnace at 1û6S°CE30min and

tempered twlce at 538'CMHRS. (Loft) 320X and (Right) 800X.

Fig: 10

Hl3 Toal steel austenitized in vacuum furnace at 102S0C/30min and

temperd twico at 538% and 593'CnHRS. (Left) 320X and (Right) 8ûûX.

and exhibited minimal signs of alloy segregation. Hardening at 1 O6S0C followed

by tempering resulted in coarsening of the martenslic structure and increased

dissolution of primary carbides as shown by the photomicrographs contained in

Fig 11. Austenitizing of tool steel at 1065°C versus 1025°C will decrease the

average ASTM grain size number by approxirnately 13%. A decrease in grain size

of 13% is not considered to be a serious degradation in the structure of a tool

steel (41). The optical micrograph also reveals that the austenitizing temperature

of 1065'~ resulted in increased delineation of the austenite grain boundaries; this

condition is caused by the presence of intergranular proeutectoid carbide. It is

suspected that the presence of the proeutectoid carbide phase in the material

austenitized at 1065 '~ followed by air cooling to room temperature will have a

deleterious effect on the totighness of this material (41).

Since increased austenitizing temperatures dissolve more carbide and since the

hardness of martensite is proportional to its carbon concentration, the effect of

austenitizing temperature on the hardness of the martensite is expected. Thus,

the trend of increasing as-quenched hardness with increasing austenitizing

temperature is related to an increase in the alloy content of the martensite.

H13 steel is hot working steel, and therefore a high austenitiang temperature

without causing grain growth is important to improve the red hardness and the

high dynamic impact value.

Fig: 11

Hl3 Tool Steel Heat treated in vacuum furnace at 1065°C/30min and temperad Nice at 538OC I2HRS. 800X

3.2 Heat treatment in open atmophere furnace

02 and Hl3 tool steel annealed samples (linxlinxlin) were heat treated at

austenitizing temperatures of 1025°C and 1038°C in an electrically heated, open

atmosphere furnace for 30 minutes, followed by air cooling. After heat treatment a

thick layer of scaling and decarborization was obsecved, which was measured

using a micro hardness tester with a Knoop diamond indentor, a 500 gram load

and a 15 second load time. A 500-gram load was used in order that the Knoop

hardness numbers could be converted accurately to HRC values. The result

(Figures 12 8 13) shows that D2 tool steel has a .35mm decarborised layer and

Hl3 tool steel has a .45mm decarborised layer. Heavy duty cutting blades and

machine knives must have sharp and thin working edges. During heat treatment

not only decarborization takes place, but also precious alloying elements that are

present in very stringent amounts in the steel are lost. Any slight loss of these

elements can be expected to reduce the tool quality below that which is predicted

(42).

3.3 Heat treatrnent by current operation

Prematurely failed heavy-duty cold and hot work machine knives samples

manufactured by 02 and Hl3 tool steels were received from Central Welding Ltd.

The steels were heat treated in an electhl ly heated, open atmosphere fumace

using an austenitizing temperature of 1 0 2 5 ~ ~ followed by tempering. The samples

were prepared very carefully in order to observe the microstructure using an

optical microscope.

The optical micrograph of 02 tool steel reveals that coarse carbide particles

(white) have a net-like microconstituent throughout the matrix (Figure 14). There

is evidence of large and nonequiaxial prior austenite grain boundaries and also

precipitation of carbides along these grain boundaries, as shown in Figure 14. The

rnartensite structure is not well revealed; we believe this is due to the presence of

retained austenite in this material. A coarse carbide network present in steel will

deteriorate the mechanical properties of the steel, and during working operation,

cracks can initiate and propagate through this brittle carbide network and fracture

can occur even if the applied stresses during working are low.

H l 3 steel heat treated at 1025°C followed by tempering is shown in Fig 15. The

micrograph reveals large prior austenite grain boundaries with clear evidence of

carbide precipitation on these grain boundaries. Hig h magnification reveals the

carbide precipitation on martensite lath boundaries (Figure 15). If during heating

the austenitizing temperature is high, the carbides will dissolve to a large extent

into solution, and grain growth will occur and the precipitation of proeutectoid

carbides on cooling will have a greater tendency to take place at coarse austenite

grain boundaries. The martensite start temperature is lower than usual in this

case and a high amount of austenite will be retained. This austenite during

working under stresses and temperature will change to upper bainite or fresh

Fig: 14

D2 Tool Steel heat treated by cunent opewation at 102S°C130min tempered

twice at 3ûû°C/2HRS. (Left) 125x and (Right) 800x

Hl3 Tool steel Heat treated In open atmosptmre furnace at 1ô2S°C and

tempered twice at 538" IZHRS.

martensite, which has a brittle structure and can deteriorate the toughness of

steel.

3.4 Charpy V-notch dynamic impact test

V-notch dynamic impact tests were carried out because in service D2 and H l3

tool steel heavy-duty machine knives and blades are used under dynamic impact

loading conditions. The objective was to determine which heat treatment condition

produced a high impact value and what is the effect of tempering temperature and

nurnber of temperings under V-notch impact loading conditions. The 02 and H l 3

tool steels samples were rnachined to precise tolerances and heat treated in an

electrically heated argon gas protective atmosphere fumace at 1025"C, 1038*C,

and 1065°C for 30 minutes. Seventy-five samples from each 02 and H l3 tool

steels were used for al1 three austenitizing temperatures. Asquenched samples of

D2 tool steel were tempered up to three times at each of 205'C, 538"C, and

593'C for two hours, and approximately twenty-five samples were used at each

tempering temperature. Asquenched samples of H l 3 tool steel were tempered at

5 3 8 * ~ , 593 '~ , and 620°c/2hrs, and approximately hventy-five samples were used

at each tempering temperature.

The results of the room temperature Charpy V-notch impact testing are plotted as

a function of austenitizing temperature, tempering temperature, and number of

temperings (Figures 16-21 ).

Heat treatment of 02 steel at 1025"C, 1038'C, and 1065°C temperatures

followed by ternpering at 5 3 8 ' ~ ~ which is within the secondary hardening range,

decrease the impact toughness values as expected (43), but the decrease is not

significant as shown in Figure 21. The impact toughness values after ternpering at

5 3 8 ' ~ were found to be lower than the impact toughness values obsewed after

tempering at 205°C and 593°C. It is speculated that the retained austenite present

in D2 steel may help to provide high impact toughness at 205'~.

Austenitizing of D2 steel at 1025'C, 1 O38"C, and 1065°C followed by tempering

up to three times at 593'C (Figure 19) shows a trend that an increase in the

number of temperings at each austenitizing temperature significantly increases

the toughness of the material. The increase in toughness between one and three

ternpers after using an austenitizing temperature of 1038°C followed by tempering

at 593'C, is 83% and between two and three tempers is 17%.

H 1 3 tool steel heat treated at the above-mentioned austenitizing temperatures

followed by tempering at 538O~, 593O~, and 6 2 0 ' ~ has a trend that impact

toughness of the steel increases with increasing the nurnber of temperings. The

results show that tempering three times versus one or two gives high impact

toughness values, as shown in Figures 16-18. The increase in toughness

between two and three tempers after using an austenitizing temperature of

1038°C followed by ternpering at 538"C, is 25% and drop in hardness is less than

1 HRC, which is not significant as shown in figure 16. The possible explanation for

the increase in toughness after the third temper is the optimum distribution of

Fig: 16 Austenite temp: 1 OZ0C, 1 O38OC, and 1 O6S0C Tempering temp: 53S°C No of tempers : 1,2, and 3

Charpy Impact Test Result Material Hl3 tool steel

1038°C 1065°C

Austenitizing temperature(%)

Fig: 17 Austenite temp: 1025OC, 1038OC, and 1065OC Tempering temp: 53fI0C, and 593OC No of tempers : 1,2, and 3

Charpy Impact Test Result Material H l 3 tool steel

1038°C 1065°C I

Austenitizing temperature('C) i / Ml One tenper at 538'CRhrs O Tw o terrpers at 538'Ct593"C12hrs B ~ h r e e tempers 538'C,593'C,593'C/2hrs /

Fig: 19 Austenite temp: 1025OC, 1 O3S0C, and 106S°C Tempering temp: 593OC No of tempers : 1,2, and 3

Charpy Impact Test Result Material 0 2 tool steel

NO of tempers at 593°C12hrs / O A ustenitized at 1025"c/30rnin O Austenitized at 1038'c/30min O Austenitized at 1065"c/30min

Fig: 20 Austenite temp: 102S°C, 1038OC, and 106S°C Tempering temp: 205OC No of tempers : 1, and 2

C harpy Impact Test Result Material D2 tool steel

No of tempers at 20S°C12hrs

I O Austenitized at 1025"C/30min O Austenitized at 1038"C/30min O A ustenitized at 1065OC130min Î

Fig: 21 Austenite temp: 1 0 2 ~ ~ C , 1038OC, and 106S°C Tempering temp: 53g0C No of tempers : 1,2, and 3

C harpy Impact Test Result Material D2 tool steel

No of tempers at 538OCIZhrs

I i O Austenitized at 1025"C/30min 13 Austenitized at 1 03B°C130min 8 Austenitized at 1065'~/30min~

alloying elements between carbides and the matrix, finer dispersion and

agglorneration of carbide particles that precipitate dunng first and second

tempering, and spherodizing of carbide particles that are present on interfacial

martensite boundaries. The increase in toughness value is also evident from the

observation made by SEM of greater ductility on the fractured surfaces after the

third tempering.

Analyses of the fracture surfaces of 02 and Hl3 steel from martensitic impact

specimens using Scanning Electron Microscopy follow.

3.5 Fracture Surface Analysis by using Scanning Electron

Microscope

The Charpy impact fractured surfaces show different morphology for the

samples tempered once, twice, and thrice at the same tempering temperature. It

is clear in the fractrograph taken at low magnification (Figure 22) that 02 tool steel

tempered once has brittle features as compared to three times tempering which

shows more plasticity, and the drop in hardness is less than PHRC which is not

significant. Due to its high carbon and high chromium content, D2 steel has

coarse chromium carbides throughout the matrix, and the fracture morphology on

the fractrograph in Figure 23 shows that fracture occun due to the breaking of

these carbide particles, and therefore D2 steel absorbs very little fracture energy.

The fractrograph in Figure 24 shows the matching part on high magnification and

reveals that the carbide particle is separated into two pieces without experiencing

Fig: 22

D2 Tool steel austenitized at 103B°C/3ûmin (Left) Tempered once at 593'C

RHRS and (Right) Temperd thrice at 593'C RHRS.

Fig: 23

D2 Tod steel austenitized at 1038OC130rnin ternpered thrice at 593'C RHRS

Showing the matching part and moiphology of fracturd surface.

Fig: 24

02 Tool steel austenitized at 1û30°CMOmin tempered thrice at 593°C DHRS

Showing the matching part of fraotured carbide.

plastic deformation, and however there is a plastically deforrned region

surrounding the carbide particle, evidenced by the appearance of many dimples.

The predominant fracture mode displayed by these particular specimens was

transgranular quasiçleavage, which is not uncornmon for hig h strength, tempered

martensitic steels (34). Figure 25 shows the morphology of the fractured surface

for a 02 steel sample austenitized at 1065°C and tempered three times at 593°C.

Even at high austenitizing temperatures, the fracture is mostly transgranular in

nature and a portion of the samples have intergranular fracture along prior

austenite grain boundaries. The transgranular nature is due to coarse carbide

particles that do not dissolve into matrix even at high austenitizing ternperatures,

as well as the agglorneration of precipitated carbides after tempering three times.

The intergranular mode of fracture is related to the segregation of carbides on

prior austenite grain boundaries.

The results of Hl 3 impact toughness test show a clear trend that toughness of

the material increases as the nurnber of temperings increases, and therefore

tempering three times after austenitizing gives higher toughness than tempering

once or twice. Austenitization of H l3 tool steel samples were carried out at

1025*C, 1038°C. and 1065'C followed by tempering up to three times at 538"C,

593'C, and 620°C. Austenitizing at 1038°C followed by tempering up to three

times gives high impact toughness values venus 1025'~ and 1065'~, whereas

austenitizing at 1065°C shows lower toughness as compared to austenitizing at

Fig: 25

D2 Tool Steel heat treated in controlled atmosphaie furnace at

1 0BS°C/30min and temperad thrice at 593OCEHRS.

1025°C. The average increase in dynamic impact toughness between two and

three temperings is 15% and an austenitizing temperature of 1038'~ followed by

tempering at 5 3 8 ' ~ resulted in an increase of 25%.

Tempering once at 538°C resulted in low impact toughness values at all three

austenitizing temperatures because of the secondary hardening that occurs in

Hl3 steel. After tempering at 53e°C, the fine particles of VC in which some

molybdenum is dissolved, precipitate throughout the rnatrix and along grain

boundaries. Also, retained austenite present in the steel transfomis into

secondary rnartensite, impairing the ductility and impact toughness of the steel

(33).

Analysis of the fracture swfaces frorn fully martensitic impact specimens

tempered at 593°C and tested at roorn temperature revealed that a distinct

change in fracture morphology occurs as the austenitizing temperature is

increased . For exarnple, the fracture surface associated with material austenitized

at 1025% indicates that considerable plastic deformation occurred during the

fracture process as is evidenced by the presence of the raised lips shown in the

scanning electron rnicrograph (Figure 26). This type of structure on a fracture

surface indicates that the material has relatively good ductility and toughness.

This statement is supported by the average Charpy V-notch impact toughness

(4.3 ft-lb) measured by the material austenitized at 1025'~ and air cooled

followed by tempering thrice ai 5 9 3 ' ~ (Figure 17). The predominant fracture mode

displayed by these particular specimens was transgranular quasi-cleavage.

However, as is evidenced by the fractrograph, a portion of the fracture surface

was intergranular in nature. Figure 27 clearly reveals that a decrease in the

transgranular quasiçleavage component and an increase in the intergranular

component of the fracture accompanies an increase in austenitizing temperature

(1 065'~). In addition, extensive secondary microcracking is readily visible in this

fractrograph. Based on the analysis of the asquenched and tempered

microstructures of the rnartensitic material, the trend of increasing intergranular

fracture with increasing austenitizing temperature is to be expected.

Fractrographs shown in Figure 26 demonstrate the difference between

specimens that were austenitized at 1025'~ followed by tempering at 5 9 3 ' ~ two

and three tirnes, respectively. The micrograph of the twice tempered specimen

shows a net-like microconstituent wRh cleavage facets, whereas tempering three

times resulted in considerable plastic deformation during the fracture process as

is evidenced by the presence of the raised lips and coarse features of the fracture

surface and by the 15% increase in impact toughness. This type of structure on a

fracture surface indicates that the material has relatively good ductifity and

toughness. The cornparison on low magnification in Figure 28 also shows the

same behavior.

From the above research, we can Say that the impact resistance of these tool

steeis is influenced by a number of physical and structural variables such as grain

size, hardness, and type and volume fraction of phases present. However, the

primary variables that were affect4 by austenitizing temperature are grain size,

Fig: 26

Hl3 Tool steel austenitized at 1025*Wmin (Left) Tempered twice at 538°C

and 593OC/WRS (Right) Tempered thrice at 538'C, 593'C and 593°C /2HRS

Fig: 27

Hl3 Tool steel austenitked at 1065*C/30min tempered thrice at 538"C,

593°C and S90°C DHRS showing the morphology of fracture surface.

Fig: 28

Hl3 Tool steel austenitized at 1 02S0~0min (Lett) Tempered Nice at 538°C

and 593OCIZHAS (Right) Tempereâ thrice at 538*C, 593'C and 593% MHRS.

hardness capability, and the type and amount of various transformation products

that are present. In general, impact resistance is inverseiy proportional to grain

size and hardness.

Austenitizing temperatures in excess of 1038'~ resulted in coarsening of the

austenitic grain structure, increased dissolution of carbides, increased tempered

hardness capability, and decreased Charpy V-notch impact toughness.

3.6 Hardness measurement in HRC

Rockwell C hardness testing was petfomed on hardened and tempered

specimens of D2 and Hl3 tool steel measured after Charpy V-notch impact

testing in order to establish a relationship between hardness of the specimen and

energy absorbed by the material. The average results are contained in tables 3.1

and 3.2. The data contained in Figures 29-34 indicate that an increase in

austenitizing temperature for a material air cooled following austenitization and

tempering resulted in increased tempered hardness. Tempering temperatures of

538"C, 593°C and 620°C used for 02 and Hl3 tool steel sarnples show similar

be havior.

Thus, the use of increased austenitizing temperatures promoted improved

temper resistance in the 02 and Hl3 tool steels that were evaluated. Undoubtedly

this effect is related to the increased levels of alloy in soiid solution that would be

available to form temper carbides. The degree of strengthening resulting from

second phase particles depends on the distribution of the particles in the ductile

TABLE 3.1- Tempered Rockwell C Hardaess and V-Notcb Charpy Impact Values for D2 steel samples (Three samples were used at each test for accuracy of results)

1 Austenitizing Temp 1 No of Temper at OC

Thrice 593 '~/2hrs

Once Twice Thrice

20s Ocnhr~

1025 o~/30min 538 "C12hrs

Once Twice

Hardness (HRC)

-

Charpy Impact Energy fi-lb

Once Twice

1 038 '~/30min 538 '~/2hrs

Once Twice

1.2 1.9 2.2

L

1

205 '~/2hrs Once Twice

593 'C/Z hrs Once Twice Thrice

58.5 56.1

48.1 47.3 46.3

I

I

1.2 1 .O

1.1 1.9 2.1

L

1065 '~/30min

t

1.1 1 .O

62.7 62.0

1.6 l

1 -9

205 '~L2hn Once Twice

538 Oc/2hrs Once Twice Thrice

593 o~/2hrs Once Twice Thrice

E

64.1 60.6

48.7 47.3 46.0

f

60.7 60.7

1.5 1.8

TABLE 3.2 - Tempered Rockuell C Hardness and V-Notch Cbarpy Impact Values for Hl3 steel simples (Tbree sampks were used ot each test for accuracy of results)

Austenitizing Temp

1025 "/30min

1 Once I I

1 Twice

1 1 Twice 1 1

I I Once 538 '~12hr I 55.8 1 2.6 I

No of Temper at OC

538 'cl2 hrs Once Twice Thrice

593 '~12hrs

50.5 1 3 .7

1 1 Twice 1 56.0 1 3 -5

Hardness (HRC)

55.8 55.0 54.2

i 1 03 8 '~/30min

--

Charpy Impact Energy fi-lb

2.6 3.4 3.9

1 1 Twice 1 44.5 1 5 -8 1

538 Ocnhr~ Once

Thrice 593 '~/2hrs

Once 538 Ocl2h.r Twice Thrice

620 '~/2hrs

56.6

55.3

56.6 50.2 49.0

2 9

3.4

2.9 4.1 4.4

2.4 3 -5

1065 o ~ / 3 ~ m i n

593 '~nhrs Once 538 '~12hr Twice

1 1 Twice 1 44.9 1 5.6 1

2.9 Once 538 Oc12hr

620 "C12hrs Once 538 '~/2hr

56.6

Once Twice

58.0 50.8

58.0 56.5

2.4 3.6

58 .O 2.4

rnatrix. For a given volume fraction of a second phase, reducing the particle sire

decreases the average distance between particles, which enhances the

precipitation strengthening effect. So during tempering, very fine carbides

precipitate throughout the matrix giving high secondary hardness.

Figures 29-34 reveal comparatively higher secondary hardness at 1 065°C

austenitizing temperature, than at 1038°C and 1025°C after tempering up to three

times at 538°C. However, tempering at 205 '~ shows the reverse behavior. A

possible explanation for this behavior is the presence of a relatively large amount

of retained austenite after air cooling from 1065'~ (versus 1038°C or 1025°C)

that, after tempering at 2 0 5 ~ ~ ~ does not transform into martensite or another

transformation product, and as a result shows relatively lower hardness. The

second or third tempering at 593 '~ or 62CI0C produced almost the same hardness

for al1 three austenitizing temperatures. This research shows t hat tempering

following austenitization at high temperature precipitates very fine, unifomly

distributed carbide particles that produce high strength and hardness, whereas

tempering following austenitization at low temperature precipitates coarser

carbide particles, which g ive relatively low strength and hardness. Second or third

tempering agglomerate the finer and coarser carbide particles in the same

manner and gives similar hardnesses.

Fig: 29 Austenite temp: 1025OC, 1038OC, and 106S°C Tempering temp: 53g0C No of tempers : 1,2, and 3

Hardness data Material D2 tool steel

1 2 3

No of tempers at 538°C12hrs

Fig: 30 Austenite temp: 1025OC, 1038OC, and 106S°C Tempering temp: 20S°C No of tempers : 1, and 2

Hardness data Material D2 tool steel

No of tempers at 20S°C12hrs

I + Austenitioed at 1 O25"CBOMN - Austenitized at 1038"C(30niin + Austenitized at 1 O65"CBOnin '

Fig: 32 Austenite temp: 103tB0C, and 106S°C Tempering tem p: 538OC, and 620°C No of tempers : 1,2, and 3

Hardness data Material H l 3 tool steel

1 2 3

No of tempers at 538OC, and 62O0CI2hrs

1 -+- Austenlizing temperature 1038°C/30min -L Austenlk ing temperature 1 065'C/30min Series3

Fig: 34 Austenite temp: 1025OC, 1038*C, and 106S°C Tempering temp: 53a°C, and 593OC No of tempers : 1,2, and 3

Hardness data Material H l 3 tool steel

-t Austenitized at 1025'Ci30nln - Austentized at 1 03B°C(30nin -*-- Austenitized at 1065'C130min 1 1

3.7 Summary of Major Results

The microstructures of Hl3 and D2 alloy steels, heat treated in an open

atmosphere furnace (as received) are coarse, with large grain sire. There is

also evidence that most of the carbides are dissolved into solution and then

precipitated along grain boundaries. This kind of microstructure would reduce

the mechanical properties considerably, especially toughness, resulting in

chipping and dulting of the tools.

The microhardness of D2 and H l3 tool steels hardened in an open

atmosphere furnace is not uniform.

Heat treatment in open atmosphere gives a thick layer of oxidation and

decarburization.

Vacuum heat treatment gives homogenous microstructure and hardness.

In both D2 and H l3 tool steels, hardening followed by three temperings gives

high impact toughness, in comparison with hardening followed by one and two

temperings.

Hardness drop after three temperings in comparison with one and two

temperings is only within 2-HRC range.

H l 3 and 02 tool steels hardened at 1038'~ followed by tempering show high

impact toughness.

D2 steel austenitized at 1038'~ followed by tempering ai 5 9 3 ' ~ shows a 17%

increase in toughness after three tempers versus two ternperings.

9. Hl3 steel austenitized at 1038'~ followed by tempering at 5 3 8 ' ~ shows a

25% increase in toughness after three temperings versus two temperings.

10. lncreasing the tempering temperature of D2 and Hl3 tool steels reduces the

hardness, except in the secondary hardening zone due to precipitation of fine

carbides throughout the martensitic matrix.

4. Conclusions

Based on the above results and discussion, the following conclusions and

recommendations can be drawn.

1. 02 and H i3 tool steels should be hardened in a controlied protective

atmosphere furnace followed by air cooling .

2. Hardening in a vacuum furnace followed by liquid nitrogen gas quenching

g ives minimum distortion, uniform martensite microstructure without

decarburization and scaling and minimum amount of retained austenite. For

the reason that quenching of 02 and Hl3 tool steels in Vacuum furnace is

moderate (slower than oil and faster than air cooling), which manage little

temperature difference between core and case of the material and results in

uniform microstructure with less distortion.

3. Three temperings are necessary for both D2 and Hl3 tool steels used for

shearing processes. Because 02 and Hl3 steels are under impact loading

conditions during shearing processes and require high strength and high

toughness, tempering third time gives higher toughness without significant loss

of strength.

4. D2 tool steel should be free from carbide segregation, which solidify during

ingot casting and form coarse and brittle networks. Since such carbides are

not greatly affected by heat treatment, heavy reduction by hot working before

die manufacturing is necessary to refine the structure.

5. D2 and Hl3 tool steels contain elements such as carbon, chromium, tungsten,

molybdenum, and vanadium in closely controlled arnounts. Any slight loss of

these elements would reduce the tool quality below the predicted values.

6. For Hl3 tool steel, decreasing the rate of quenching leads to a gradua1

increase in the width of the bainite laths together with an increase in volume

fraction of upper bainite, and this results in deterioration of the toughness.

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