The Microstructure, Hardness, Impact

THE MICROSTRUCTURE, HARDNESS,

IMPACT TOUGHNESS, TENSILE DEFORMATION AND

FINAL FRACTURE BEHAVIOR OF FOUR SPECIALTY HIGH STRENGTH STEELS

A Thesis

Presented to

The Graduate Faculty of The University of Akron

In Partial Fulfillment

of the Requirements for the Degree

Master of Science

Manigandan Kannan

August, 2011 THE MICROSTRUCTURE, HARDNESS,

IMPACT TOUGHNESS, TENSILE DEFORMATION AND

FINAL FRACTURE BEHAVIOR OF FOUR SPECIALTY HIGH STRENGTH STEELS

Manigandan Kannan

Thesis

Approved: Accepted:

______Advisor Department Chair Dr. T.S. Srivatsan Dr. Celal Batur

______Faculty Reader Dean of the College Dr. C.C. Menzemer Dr. George.K. Haritos

______Faculty Reader Dean of the Graduate School Dr. G. Morscher Dr. George R. Newkome

______Date

ABSTRACT

The history of steel dates back to the 17th century and has been instrumental in the betterment of every aspect of our lives ever since, from the pin that holds the paper together to the automobile that takes us to our destination steel touch everyone every day.

Pathbreaking improvements in manufacturing techniques, access to advanced machinery and understanding of factors like heat treatment and corrosion resistance have aided in the advancement in the properties of steel in the last few years. This thesis report will attempt to elaborate upon the specific influence of composition, microstructure, and secondary processing techniques on both the static (uni-axial tensile) and dynamic

(impact) properties of the four high strength steels AerMet®100, PremoMetTM290, 300M and TenaxTM 310. The steels were manufactured and marketed for commercial use by

CARPENTER TECHNOLOGY, Inc (Reading, PA, USA). The specific heat treatment given to the candidate steels determines its microstructure and resultant mechanical properties spanning both static and dynamic. Test specimens of the steels were precision machined and conformed to standards specified and prescribed by the American Society for Testing Materials (ASTM) for both tensile tests and Charpy V-Notch impact tests.

Based on similarity of the secondary processing technique the candidate specialty steels were divided into two groups: (i) AerMet®100 and PremoMetTM290, (ii) 300M and

TenaxTM310. The impact toughness response and resultant fracture behavior of the steels were studied at different temperatures ranging from -180°C to +170°C. Tensile tests

iii were performed at room temperature and the final fracture behavior of the candidate steel was established at both the macroscopic and fine microscopic levels. The intrinsic microscopic mechanisms governing the impact toughness, quasi static deformation and final fracture behavior of each of the chosen high strength steels will be elaborated upon in light of the conjoint and mutually interactive influences of composition, intrinsic microstructural effects, and nature of loading.

ACKNOWLEDGEMENTS

First and foremost, I would like to thank my advisor Dr. T.S. Srivatsan. He was instrumental in every move I ever made in this university and I would be grateful all my life for the support rendered to me during my small stint in this university. I would like to thank Dr Celal Batur for awarding me with a Teaching assistant which helped me to complete my Master of Science degree in the Department of Mechanical Engineering. I would like to extend my sincere thanks to Dr. Craig Menzemer and Dr. Gregory

Morscher for serving on my thesis committee.

I would like to thank Michael Schmidt from Carpenter Technology,INC located at

Reading Pennsylvania for providing the materials that made this research possible.

Furthermore, I would like to give sincere thanks and recognition to the following individuals for their contribution in this research by way of knowledge and assistance:

(i) Mr. Nurudeen Balogun and Mr. Udaykar Bathini (former graduate students)

for assisting me during the initial stages of testing.

(ii) Dr. Sergei F. Lyuksyutov(Associate professor in the department of physics)

for his help with the atomic force microscope.

(iii) Ms. Michelle Petraroli (Senior Research Technician, of the Timken Company)

for assistance in microstructural observations

(iv) Mr. Clifford Bailey (Senior Engineering Technician, Department of

Mechanical Engineering), for instruction and timely assistance with intricacies

related to use of Mechanical Test Machine.

(v) Mr. Thomas J. Quick (Research Associate, of the Department of Geology), for

assistance with the use of the Scanning Electron Microscope.

(vi) Mr. Stephen Gerbetz (Senior Engineering Technician, Department of

Mechanical Engineering),

(vii) Mr. Dale Ertley (Senior Engineering Technician) and Mr. Bill (Engineering

Technician), of the College of Engineering Machine shop for assistance with

the preparation of samples for scanning electron microscopy and optical

microscopy.

(viii) Mr. David McVaney (Senior Engineering Technician), for allowing me to use

the hardness tester in the Civil Engineering Department.

Above all I want to extend my gratitude to my parents, my brother and my dear friends for all their love and support throughout my life.

TABLE OF CONTENTS

Page

LIST OF TABLES ix

LIST OF FIGURES x

CHAPTER

I. Introduction 1

II. Review of the Literature 6

2.1 Processing and microstructure of Specialty steels 6

2.2 The impact fracture toughness of high strength 10

2.3 The tensile properties of high strength steels 13

III. The Test Materials 17

IV. Test Sample Preparation

4.1 Impact tests 21

4.2 Tensile tests 21

V. Experimental Procedures

5.1 Initial microstructure characterization 24

5.2 Characterization of Microstructure using Atomic Force 24 Microscopy (AFM).

5.3 Hardness testing 25

5.3.1 Macro-hardness tests 26

5.3.2 Micro-hardness tests 28

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5.4 Impact tests 30

5.5 Mechanical tests (tensile) 31

5.6 Failure-damage analysis 31

VI. Results and Discussion

6.1 Initial Microstructure 33

6.2 Initial Microstructure using the Atomic Force Microscope. 34

6.3 Hardness

6.3.1 Microhardness measurement 44

6.3.2 Macrohardness measurement 45

6.4 Impact Toughness 51

6.5 Dynamic Fracture 55

6.6 Tensile Properties 81

6.7 Tensile fracture behavior 89

VII. Conclusions

7.1 Impact Toughness 97

7.2 Tensile Deformation and Fracture Analysis 98

References 100

Appendix 105

APPENDIX A PROCEDURE FOR PERFORMING 105 THE TENSION TEST ON THE INSTRON-8500 SERVO HYDRAULIC TESTING MACHINE

APPENDIX B PROCEDURE FOR PERFORMING LOOP 109 SHAPING FOR A GIVEN STRESS RATIO AND MATERIAL ON THE INSTRON-8500 SERVO HYDRAULIC TESTING MACHINE

viii

LIST OF TABLES

Table Page

3.1 Nominal chemical composition of the chosen materials ………………..19

3.2 Primary and Secondary Processing history on the chosen ……………...20 High strength steels

5. 1 A compilation of the micro-hardness measurements made on the Specialty steels samples…………………………………………………26

5. 2 A compilation of the macrohardness measurements made on the Specialty steels samples...... 28

6. 1 The impact toughness properties of the four high strength steels in N-m….………………………………………………………………..55

6. 2 A compilation of the room temperature tensile properties of the test Materials………………………………………………………………...84

6. 3 A summary of the monotonic mechanical parameters of the candidate.. 86

LIST OF FIGURES

Figure Page

4.1. A schematic showing dimensions of the Charpy V-Notch test specimen………22

4.2. A schematic of the cylindrical test specimen used for …………………………23 Mechanical testing(Tensile).

6.1. Optical micrographs showing the key micro-constituents of AerMet® 100 at three different magnifications ...... 36

6.2. Optical micrographs showing the key micro-constituents of PremoMetTM 290 at three different magnifications ...... 37

6.3. Optical micrographs showing the key micro-constituents of 300M at three different magnifications ...... 38

6.4. Optical micrographs showing the key micro-constituents of Tenax TM 310 at three different magnifications ...... 39

6.5. Atomic force microscope image of the etched surface of AerMet® 100 showing profile and nature of roughness on a 10 μm x 10 μm section ...... 40

6.6. Atomic force microscope image of the etched surface of PremoMetTM 290 showing profile of roughness on a 10 μm x 10 μm section ...... 41

6.7. Atomic force microscope image of the etched surface of 300M showing nature and profile of the roughness on a 10 μm x 10 μm section ...... 42

6.8. Atomic force microscope image of the etched surface of TenaxTM 310 showing nature and profile of roughness on a 10 μm x 10 μm section ...... 43

6.9. A profile showing the variation of microhardness and macrohardness value across the length of AerMet® 100...... 46

6.10. A profile showing the variation of microhardness and macrohardness values across the length of PremoMetTM 290 ...... 47

6.11. A profile showing the variation of microhardness and macrohardness values across the length of 300M ...... 48

6.12. A profile showing the variation of microhardness and macrohardness values across the length of TenaxTM 310 ...... 49

6.13. Bar graph depicting the average microhardness and macrohardness values of the four specialty steels...... 50

6.14. Variation of energy absorbed (N-m) as a function of test temperature for the two steels AerMet® 100 and PremoMetTM290 ...... 53

6.15. Variation of energy absorbed (N-m) as a function of test temperature for the two steels 300M and TenaxTM 310 ...... 53

6.16. Bar graph comparing the energy absorbed with test temperature of the four chosen materials:AerMet® 100, 300M, PremoMetTM 290 and TenaxTM 310 ...... 54

6.17. Scanning electron micrographs of Aermet® 100 deformed under impact following exposure in liquid nitrogen at – 180°C, showing: ...... 58

(a) Overall morphology of failure. (b) High magnification of (a) showing a healthy population of microscopic voids and dimples. (c) Morphology of the fine microscopic voids and their distribution (d) Isolated fine microscopic cracks reminiscent of locally brittle failure mechanisms.

6.18. Scanning electron micrographs of AerMet® 100 deformed under impact following exposure in dry ice at – 55°C, showing: ...... 59

(a) Overall morphology of failure. (b) High magnification of (a) showing a random distribution of fine microscopic cracks. (c) High magnification of (b) showing a random distribution of fine microscopic voids and a healthy population of ductile dimples. (d) The size and shape of the dimples on the overload surface

6.19. Scanning electron micrographs of Aermet® 100 deformed under impact in the room temperature environment at 25°C, showing: ...... 60

(a) The as-ruptured morphology of failure (b) High magnification observation of the region immediately adjacent to the notch. (c) Macroscopic crack surrounded by a healthy population of shallow dimples. (d) Fine microscopic cracks separating the region in the vicinity of notch and overload

6.20. Scanning electron micrographs of Aermet 100 deformed under impact following exposure in a furnace at 170°C,showing…………………...…………61

(a) Overall morphology of failure. (b) High magnification observation in the region immediately adjacent to the notch showing isolated fine microscopic cracks and a population of dimples (c) High magnification observation of (b) showing the size and Morphology of the microscopic cracks. (d) Microscopic voids and pockets of shallow dimples.

6.21. Scanning electron micrographs of PremoMetTM 290 deformed under impact following exposure in liquid nitrogen at – 180°C, showing: ...... 65

(a) Overall morphology of failure as damage radiates outward from the notch. (b) The region adjacent to the notch showing a healthy population of fine microscopic voids. (c) High magnification of (b) showing the morphology and shape of the microscopic voids. (d) An array of fine microscopic cracks on the overload fracture surface.

6.22. Scanning electron micrographs of PremoMetTM 290 deformed under impact following exposure in dry ice at – 55°C, showing: ...... 66

(a) Overall morphology of the fracture surface. (b) High magnification observation in the region immediately adjacent to the notch showing a random distribution of microscopic voids. (c) Microscopic crack reminiscent of locally brittle failure. (d) A mixture of both macroscopic and fine microscopic voids in the region away from the tip of the notch

6.23. Scanning electron micrographs of PremoMetTM 290 deformed under impact in the room temperature (25°C) laboratory air environment,showing ...... 67

(a) Overall morphology of failure (b) The region adjacent to the notch showing a healthy population of fine microscopic voids and isolated microscopic cracks. (c) High magnification of (b) showing the morphology, size and distribution of the microscopic voids. (d) High magnification observation of ( c) showing shallow

xii

nature and distribution of the dimples intermingled with fine microscopic voids.

6.24. Scanning electron micrographs of PremoMetTM 290 deformed under impact following exposure in a furnace at 170°C, showing: ……………..68

(a) Overall morphology of the impact fracture surface. (b) The region adjacent to the notch showing features reminiscent of ductile and brittle failure mechanisms. (c) An array of fine microscopic cracks intermingled with microscopic voids. (d) The size, shape and shallow nature of the dimples in the region away from the notch tip.

6.25. Scanning electron micrographs of 300M deformed under impact following exposure in liquid nitrogen at – 180°C, showing: ...... 71

(a) Overall morphology of failure. (b) High magnification observation in the region immediately adjacent to the notch showing voids of varying size and their random distribution. (c) High magnification observation of the impact fracture surface away from the notch. (d) A healthy population of shallow dimples of varying size and shape.

6.26. Scanning electron micrographs of 300M deformed under impact following exposure in dry ice at – 55°C, showing: ………...72

(a) The overall morphology of the impact fracture surface. (b) High magnification observation in the region adjacent to the notch showing a population of voids of varying size and healthy distribution of dimples. (c) High magnification of (b) showing the morphology and shape of microscopic void and dimples. (d) The shallow nature, and non-uniform size and shape of the dimples in the region away from the notch.

6.27. Scanning electron micrographs of 300M deformed under impact in the room temperature (25°C) laboratory air environment, showing: ...... 73

(a) The overall morphology of failure as it radiates away from the notch. (b) High magnification observation in the notch region showing a random distribution of macroscopic and fine microscopic voids.

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(c) The distribution and shape of shallow dimples and fine microscopic voids. (d) Isolated microscopic crack in a predominantly ductile region.

6.28. Scanning electron micrographs of 300M deformed under impact following exposure in a furnace at 170°C, showing ...... 74

(a) Overall morphology of failure as it radiates outward from the root of the notch. (b) High magnification observation in the region adjacent to the notch reminiscent of locally ductile failure mechanisms. ( c) Microvoid coalescence to form fine microscopic crack surrounded by dimples (d) Microcracks and random distribution of shallow dimples in the region of overload.

6.29. Scanning electron micrographs of TenaxTM 310 deformed under impact following exposure in liquid nitrogen at – 180°C, showing: ...... 77

(a) Overall morphology of the fracture surface. (b) High magnification in the region adjacent to the notch tip. (c) High magnification of (b) showing a healthy population of microscopic voids, dimples and microscopic cracks. (d) Shallow nature of the dimples and coalescence of microscopic voids to form microscopic cracks.

6.30. Scanning electron micrographs of TenaxTM 310 deformed under impact following exposure in dry ice at – 55°C, showing: ...... 78

(a) Overall morphology of the fracture surface (b) High magnification in the region adjacent to the notch showing voids of varying size and isolated microscopic cracks. (c) High magnification of (b) showing growth and eventual

coalescence of microscopic cracks to form macroscopic crack surrounded by a population of dimples. (d) The size, shape and distribution of dimples.

6.31. Scanning electron micrographs of TenaxTM 310 deformed under impact in the room temperature (25°C) laboratory air environment, showing: ...... 79

(a) Overall morphology of the fracture as it radiates outward from the notch tip. (b) High magnification showing the features adjacent to the notch. (c) High magnification of (b) showing the size and distribution of

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the microscopic voids. (d) Macroscopic cracks surrounded by population of fine microscopic voids and dimples.

6.32. Scanning electron micrographs of TenaxTM 310 deformed under impact following exposure in a furnace at 170°C, showing: ...... 80

(a) Overall morphology of the impact fracture surface. (b) The region immediately adjacent to the notch showing a population of microscopic cracks, microvoids and shallow dimples. (c) High magnification of (b) showing profile of a typical microscopic cracks. (d) Distinct macroscopic crack in the region away from the notch tip.

6.33. Bar graph depicting the elastic modulus of the four candidate high strength steels...... 83

6.34. Bar graph comparing the Yield Strength (MPa) and Ultimate TensileStrength (MPa) of the four chosen and studied …………………… ….84

6.35. The engineering stress versus engineering strain curve for the four high strength steels deformed in tension at room temperature (T = 25°C). ……..85

6.36. Variation of true stress (MPa) with true strain (%) for AerMet® 100 deformed in tension at room temperature (T=25°C). ………………87

6.37. Variation of true stress (MPa) with true strain (%) for PremoMet™ 290 deformed in tension at room temperature (T = 25oC)...... 87

6.38. Variation of true stress (MPa) with true strain (%) for 300M deformed in tension at room temperature (T = 25oC)...... 88

6.39. Variation of true stress (MPa) with true strain (%) for Tenax™310 deformed in tension at room temperature (T = 25oC)...... 88

6.40. Scanning electron micrographs showing tensile fracture surface of AerMet®100 deformed at room temperature (25° C), showing: ...... 93

(a) Overall morphology of failure. (b) High magnification of (a) showing a healthy population of voids of varying size inter-dispersed with a population of dimples xv

(c) High magnification observation of the region of overload showing microscopic voids and microscopic cracks. (d) The shallow nature of dimples and microvoid coalescence to form microscopic cracks.

6.41. Scanning electron micrographs of the tensile fracture surface of PremoMetTM290 deformed at room temperature (T = 25°C), showing: ...... 94

(a) Overall morphology of failure (b) Near featureless transgranular region. (c) High magnification of (b) showing a population of fine microscopic cracks intermingled with microscopic voids. (d) The region of overload covered with shallow dimples and population of microscopic voids of varying size and shape.

6.42. Scanning electron micrographs of the tensile fracture surface of 300M deformed at room temperature (T = 25°C), showing: ...... 95

(a) Overall morphology of failure (b) High magnification of (a) showing a healthy population of microscopic voids intermingled with dimples. (c) High magnification observation in the region of overload showing cracking at grain boundary triple junction. (d) Grains covered with shallow dimples and microscopic voids, reminiscent of locally ductile failure mechanisms.

6.43. Scanning electron micrographs of the tensile fracture surface of ...... 96 Tenax TM 310 deformed at room temperature (T = 25°C), showing:

(a) High magnification observation of the transgranular region showing fine microscopic voids of varying size intermingled with fine microscopic cracks. (b) A section of the transgranular fracture surface showing a distribution of fine microscopic cracks of varying size and shape. (c) The region of tearing or overload showing voids of varying size intermingled with pockets of ductile dimples.

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CHAPTER I

INTRODUCTION

The story of alloy steels may have its beginning before the dawn of history.

However the last fifty years have witnessed more progress in the alloying of special steels and putting them to work effectively than the millenniums preceding. The search for newer materials that can offer improved properties over the existing counterparts is a never ending process. This is because of the challenging requirements for new and improved materials with the evolution of novel processing techniques and other technologies. One way of achieving this goal is by choosing the right alloying elements and the advantages of adding alloying elements are as follows [1].

A. An improvement in mechanical properties

- Improved strength

- Increase in toughness and ductility.

- Improvement of mechanical properties at a wide range of temperatures.

B. An improvement of chemical inertness.

- Protection against corrosion.

- Stability at elevated temperatures.

Steel is an alloy of mostly iron and carbon and the amount of carbon is between

0.1% to 2.1% anything above 2.1% in an alloy changes the steel into cast iron [2]. Based

1 on the carbon percentage the steels are divided into low carbon steels, medium carbon steels and high carbon steels. Alloying elements are the hardening agent and helps in preventing dislocations in the iron lattice [1]. The amount of alloying elements added to the steel controls the hardness, strength and ductility of the material. The alloying elements, which can be added to steels are chromium, cobalt, manganese, molybdenum, nickel, phosphorus, titanium, tungsten and vanadium. The increase in strength of steels is attributed to the type of processing done to the material and the resultant microstructure.

Martensitic steels have the highest strength and this transformation is achieved by rapid cooling resulting due to quenching in water or oil and thus forming either lath or Plate morphology [3]. Increase in hardness and strength has a direct influence on ductility. As the strength increases the ductility decreases.

A few of the characteristics desired in a high strength material for possible selection and potential use for both performance-critical and non-performance-critical applications include enhanced durability, a high strength-to-weight (σ/ρ) ratio, adequate damage tolerance capability, acceptable efficiency, reliability, relative ease of manufacturing and overall cost effectiveness [3-5].A viable way to effectively compromise strength over ductility or building a perfect balance between the two is what most manufacturers of metals and alloys have been trying to achieve through the years, i.e., since the early 1970. The steels chosen for this particular study are no exception.

High strength steels often have a metallic coating so as to protect it from corrosion [6]. Thereby increasing the risk of failure due to hydrogen embrittlement and this can occur in different ways. This must be suppressed and failure must not occur.

Regardless of the material chosen, several factors are intrinsically involved in ensuring

2 the safety, reliability and efficiency of both structural and mechanical systems. The factors include sound design, proper material selection, application of economically feasible and technically sound fabrication procedures and varied construction practices

[7]. In almost every case, fracture is the final stage of any component and occurs due to factors, which are independent of the combined influence of improved design, incorrect material selection, incorrect fabrication procedures, flaws present in the material or incorrect construction practice, and the environment place a very pivotal role.

Development of newer products and finding appropriate applications for high strength steel necessitates that a careful consideration be given to all pertinent limit states, design procedures, fabrication practices and construction processes combined with a careful evaluation of all potential end uses so as to help ensure safe and reliable operation.

The mechanical properties to include fracture toughness of high strength steel are often governed by the independent or conjoint influences of chemical composition, processing history and development of intrinsic microstructural features, geometry or part thickness, temperature of operation, loading-rate, and presence of constraints at the crack (flaw) tip

[8-14]. Conditions that tend to minimize constraints in a structure are fairly beneficial for enhancing the ability of a metal to plastically deform. This facilitates in improving the fracture toughness and overall damage tolerance capability of the high strength steel.

Factors such as an increase in thickness of the part, incorporation of engineering notches, an increase in the loading rate; and lowering of the operating temperature are known to be detrimental to both ductility and fracture toughness of the chosen metal and related engineering structure [15-20].

High strength steels can be considered as a candidate for use in applications such as manufacturing of armors as it has high toughness and hardness and can withstand impact loads. It can also be used for the manufacture of fasteners, and in manufacture of landing gears in the aerospace industry as it has a high load carrying capacity. It can be used in making jet engine shafts as the load carrying capacity of these steels are high at high temperatures. Because of its high load carrying capabilities it is of great use in structural members and it can reduce the weight of structures as the thickness of structures can be drastically reduced. This helps in reducing the money spent on manufacturing and other heat treatment processes. High strength steels have great application in the automobile industry as it could be used in the manufacturing of frames and engine shafts [21-22]. These steels could be used in power industry as it could be used in manufacturing of structural tubing for ultra super critical boilers [23].

In this thesis document is presented and discussed the results of a study aimed at studying the influence of alloy chemistry, processing and test temperature on impact response, tensile response and fracture behavior of the four high strength steels, and comparing their behavior under identical conditions. The microstructure of the as- received steel was examined and characterized for the nature and morphology of the grain structure and distribution of other intrinsic features in the microstructure. The tensile tests were done on a fully automated machine at room temperature. The impact tested and failed samples of high strength steels was then examined in a scanning electron microscope for understanding the impact fracture behavior, especially the role of test temperature and contributions from microstructural features in governing the kinetics of fracture. The tensile fracture of each steel and the various factors contributing to failure

4 are briefly discussed in light of the conjoint and mutually interactive influences of intrinsic microstructural features, nature of loading and stress (load)-deformation- microstructural interactions.

CHAPTER 2

REVIEW OF THE LITERATURE

2.1 Processing and microstructure of high strength steel

Processing and microstructure plays a very important role in increasing the strength of steels. Factors such as grain size, grain structure and processing variables are very important. These steels generally have a low carbon content (0.2%-0.4%) and the strength is reached by the formation of martensite in most cases [24-27]. The austenite in carbon steels strongly influences the transformation and deformation behavior of the heat treated steels. Therefore if the austenite grain size in steel is coarse, fewer nucleation sites are available and the nucleation of the grain is reduced or slowed, thereby increasing the hardenability [28]. Austenite grain size also affects martensite transformation, in the iron- nickel (Fe-Ni) and iron-nickel-carbon ( Fe-Ni-C) alloys. Decreasing austenite grain is attributed to the higher strength of fine grained austenite, which in turn increases the shear resistance of the austenite to-martensite transformation. The reduction of austenite grain size in low carbon steels therefore offers the possibility of a significant improvement in properties of the ferritic microstructures, and plays an important role in the high strength low alloy (HSLA) steels [28].

In hardened steels, formation of martensite in fine grained austenite is preferred because of the improved mechanical properties that results. The increase in yield or flow strength of the low carbon lath martensite is unique. This is because, the martensite units

form in parallel arrays called packets [28-31].

Martensite has long been used to designate the hard microstructure found in quenched carbon steel [32]. Martensitic transformation also occurs in many nonferrous systems such as, copper-aluminum (Cu-Al) and gold-cadmium (Au-Cd) alloys [33]. In

Fe-C alloys and steels austenite is the parent phase that transforms to martensite on cooling. The martensitic transformation is diffusionless and the martensite has exactly the same composition as the parent austenite and this is upto 2% C [32]. Since diffusion is suppressed usually by rapid cooking the carbon atoms do not partition themselves between cementite and ferrite but instead are trapped in the octahedral sites of a body- centered cubic (BCC) structure thus producing a new phase called martensite [28].

Martensite is a unique phase that forms in steels. It has its own crystal structure and composition and is separated by well defined interfaces from other phases, further it is a metastable phase this is present because diffusion has been suppressed. Two major morphologies of the martensite are lath and the plate, and these develop in heat treatable steels. The ‘lath’ designation is used to describe the broad shaped units of martensite that form in low and medium carbon steels while the plate designation accurately describes the shape of the martensite units that form in the high carbon steels. The terms ‘lath’ and

‘plate’ referring to the 3-D shapes of the individual martensite laths or plates are often revealed by polishing and etching [28, 34].

Plate martensite was often not emphasized in the literature until electron microscope was invented. The units of plate martensite are well within the size range resolvable in the light microscope and frequently the retained austenite that coexists with the martensite in high carbon alloys helps in defining the plates in the light microscope.

Although the plate martensites are important in some heat treated applications, however most hardenable steels often have low or medium carbon content and as a result the lath martensites has industrial significance. Tool steels have plate martensite microstructures.

A martensite microstructure is the hardest microstructure that can be produced in any steel and results only when the transformation of austenite to a combination of ferrite and cementite is avoided [34].

Martensite, the object of the quenching treatments is quite hard but with hardness results brittleness. The brittleness of the martensite microstructures is due to a number of factors that may include the following, (i) lattice distortion caused by carbon atoms trapped in the octahedral sites of the martensite, (ii) impurity atom segregation at austenite grain boundaries, (iii) carbide formation during quenching, and (iv) residual stresses produced during quenching.

Tempering is the heat treatment of hardened steels so as to achieve a reduction of brittleness or increased toughness as the major objective. Ultimately, it is a balance of hardness and toughness required in service that often determines the conditions of tempering for a given application [28].

In addition to increasing hardenability certain alloying elements help in retarding the rate of softening during tempering. The most effective elements in this regard are the strong carbide formers, such as, chromium, molybdenum, and vanadium. Without these elements the iron-carbon alloys and low-carbon steels tend to soften rapidly with increasing tempering temperatures. This softening is largely due to the rapid coarsening of cementite with increasing tempering temperature, a process that is dependent on the diffusion of carbon and iron.

The processing procedures to the strengthening and composition of a materials. In this research study the steels my test study, the steels were processed by

(i) Vacuum induction melting (VIM) and

(ii) Vacuum arc-re-melting (VAR).

Vacuum induction melting (VIM) is one of the most common processes in secondary processing of metals. It makes possible effective degassing of the melt and precise adjustment of the alloy composition. The application of vacuum in the induction melting process is indispensable for the production of high purity metals that tend to react with atmospheric oxygen. The vacuum melting process limits the formation of non-metallic oxide inclusions that are responsible for premature failure of the structure or component.

Vacuum provides an extremely low gas content in the melt and avoids oxidation of sensitive alloy elements. Backfilling with inert gas after melting guarantees purity of the

melt, which is a result of the vacuum treatment to be preserved [36,37].

Advantages of VIM are it enables an extremely precise adjustment of the alloy composition and melt homogenization since melt temperature, vacuum, gas atmosphere, pressure and kinetics can be adjusted independently and the impurities can easily be removed [37].

Vacuum-arc re-melting is widely used to improve the cleanliness and refine the structure of standard air-melted or vacuum induction melted ingots. VAR steels and superalloys as well as titanium and zirconium and its alloys are used in a great number of high-integrity applications where cleanliness, homogeneity, improved ductility and fracture toughness of the final product are essential. The solidification structure of a VAR ingot of a material is a function of the local solidification rate and the temperature

gradient at the liquid/solid interface. To achieve a directed dendritic primary structure, a relatively high temperature gradient at the solidification front must be maintained during the entire re-melting process. The growth direction of the cellular dendrites conforms with the direction of the temperature gradient, i.e., the direction of the heat flow at that moment [38]. The primary benefits of VAR are:

(a) Removal of dissolved gases, such as hydrogen, nitrogen and carbon monoxide.

(b) A reduction of undesired trace elements with high vapor pressure.

(d) The achievement of directional solidification of the ingot from bottom to top, thus

avoiding macro-segregation and reducing micro-segregation [38].

2.2 The Impact fracture toughness of high strength steels.

Speich and Spitzig have investigated the change in ductility and toughness on inclusion of manganese sulfide in 4340 plate steels having varied sulfur content and hardened to different strengths [39]. Sulfide inclusions act plastic when hot rolled and tend to form flattened elongated shapes and decrease the ductile fracture resistance. As a result, the tensile ductility and Charpy V-notch impact energy decreased more rapidly with an increasing volume fraction of inclusions in the through-thickness plate orientation than in the transverse plate orientation [39].

According to G Krauss [40], microalloying additions of vanadium to medium- carbon steels effectively increase the strength of forged steels without subsequent heat treatment after cooling of the forgings. Enhanced strengthening is accomplished by the precipitation of fine vanadium carbon-nitride particles in the ferrite of the direct cooled ferrite-pearlite microstructures. As a result, the resistance to cleavage fracture during

impact loading was found to be low. Toughness is increased by (a) lowering the carbon content and reducing the amount of pearlite in the microstructure, (b) titanium additions to refine austenite grain size and pearlite colony size, and by increasing the sulfur content, which in vanadium-containing steels stimulates the formation of intra-granular ferrite and effectively reduces the pearlite colony size.

According to Li Jie and co-workers [41] performing a study on the influence of inclusions and voids on the fracture toughness of AerMet 100 and was compared to

AF1410 steels. And the samples of AerMet 100 were normalized at 900°C for 1 hour, air cooled (AC), reheated to 680°C, held for 8 hours, then austenitized at 885°C for 1 hour, oil quenched (OQ), cryogenic treated at -73°C for 1 hour, air warmed (AW), and aged at the temperatures of 454°C -510 °C separately for 5 hours. The following were the observed findings

(i) When aging in 454°C -510"C, ultimate tensile strength decreases with increasing

aging temperature, yield strength increases and then drops, the difference

between the two strengths reduces, while impact and fracture toughness increase

with an increasing temperature.

(ii) The fracture toughness of AerMet100 steel is influenced by microstructures and

inclusions greater than that of the AF1410 steel.

The effect of zirconium additions on the impact toughness of the heat affected zone (HAZ) in a high strength low alloy pipeline steel was studied by A.M Guo and co- workers [42]. According to them ZrO2 is formed in the liquid steel when zirconium is added during the steel making process. MnS precipitates on pre-formed ZrO2 on solidification, due to their structural similarities, which reduce the interfacial energy. As a

result, in the Zr-containing steel, the MnS inclusions tend to transform from the irregular forms, seen in Zr-free steels to small spherical complex inclusion particles. In addition, the precipitation of manganese sulphide (MnS) on the pre-formed zirconium oxide (ZrO2) facilitates the formation of intragranular ferrite. The net effect of these actions is the presence of smaller and more uniform dimpled features on the fracture surfaces of HAZ specimens of the Zr-containing simulated welds. The impact toughness of the HAZ regions has been improved by the addition of zirconium to steel.

R.Cao and co-workers [43] investigated the difference between coarse grain and fine grain 18-Nickel-200 grade maraging steel and found the toughness of the coarse grain steel(100-150μm) to be higher than that of fine grain(30-50μm) steel. The superior toughness of coarse grain-specimens than that of fine grain specimens appears only in notched specimen due to its high stress triaxiality, where the specimen is fractured at a lower fracture strain. This strain level is lower than that needed to nucleate voids by rupture of the tiny carbide plates precipitated inside the bainitic laths or the bainite laths in coarse grain specimens.

Fairchild and co-workers [44] investigated the mechanism of brittle fracture in microalloyed steel. They developed a mechanistic model to explain the results [45].

Cracked Charpy samples were used to determine the toughness of steels exposed to gleeble programs to attain various microstructures affected by heat and to simulate the heat-affected zones (HAZ). The heat-affected zone is the initiator of cleavage fracture.

Titanium nitride (TiN) inclusions present in the material was found to be the reason for cleavage fracture, thus affecting the fracture toughness of the chosen steel and not the heat-affected zones.

The Charpy V notch (CVN) test is effective in determining the impact properties in both a time effective and cost efficient manner. Samples are easily fabricated, requires a minimum amount of material and testing is quick and easy. The CVN tests are thus ideal for quality control and provide quality dynamic toughness properties while other mechanical tests offer properties determined under, condition of quasi- static loading.

2.3 The Tensile Properties of High Strength/Specialty Steels.

The mechanical properties of a material, which are functions of the microstructure, can be determined by a tension test [46]. Few researchers have developed, tested and published on various high strength steels bringing out the distinct microstructural features and their properties. Following is some of the work done by the prominent researchers in the field.

WS Du and co-workers [47] investigated the fracture behavior of 9% nickel 1000

(MPa) high-strength steel at various temperatures. They found the tensile properties to be almost the same for test specimens taken from the three different orientations

(longitudinal, transverse and through the thickness (short-transverse)). With test temperature decreasing, yield strength, ultimate tensile strength and the true fracture stress showed an increasing trend. Four critical stresses were found for this material. The fracture behavior of the specimens taken from the rolling direction specimens at various temperatures were determined by their temperature-dependent values. At −196°C, these specimens are composed of boundaries of austenite grains, a longitudinal crack initiated at the center of the neck region and propagated along the tensile stress direction to regions close to both ends of the necking. The crack then changes the orientation and develops into two transverse cracks, which propagated in two opposite directions on two

cross-sections. In the range of −150°C to −60°C for rolling direction specimens taken from the transverse crack initiated at the center of the necked region. The shear stresses produced by the transverse crack and by the necking initiated and propagated two shear cracks.

T.S Srivatsan and co-workers[48] published their research findings on the influence of the microstructure, tensile response and resultant fracture behavior characteristics of a tool-steel reinforced with titanium carbide (TiC) particles and a tool steel having no reinforcement and the following observations were noted. The microstructure of the tool steel/TiC composite revealed the reinforcing titanium carbide particles to be distributed uniformly through the metal matrix.

The elastic modulus of the composite was 10% more than the elastic modulus of the matrix alloy with no reinforcement. An increase in test temperature was found to have little influence on elastic modulus. At ambient-temperature the yield strength increased with an increase in reinforcement content in the tool-steel matrix. With an increase in test temperature the yield strength of the material degraded for the composite

. The degradation in yield strength was as high as 20%, while the ultimate tensile strength degraded only 5% . The increased strength of the TiC-reinforced tool-steel metal matrix is rationalized in terms of the concurrent and mutually interactive influences of the following: (i) work hardening arising from the generation of geometrically necessary dislocations, (ii) constrained plastic flow and triaxiality in the tool-steel matrix, and (iii) the strengthening arising from the substructure that evolves from the presence of additional dislocations.

In the study by F Ozturk [49] and co-workers, an advanced high strength steel

[denoted as DP600] was investigated in their research and following findings. The elongation obtained in an uniaxial tensile test provided the decrease with increase in temperature except in the transverse direction. The elongation was approximately 22% and there was no noticeable difference in elongation of the specimens tested at room temperature (25°C) and 200°C .The tensile strength improved when tested above 200°C and this is due to strain hardening of the material. For a given temperature and strain rate, the strain hardening coefficient was fluctuating and decreased at 200°C .Overall, the material showed a complex behavior for different temperatures and test direction.

The mechanical properties and microstructures of AISI 4340 high strength alloy steel was investigated under various tempering condition by Woei-Shyan Le and co- workers [50]. They found the strength, hardness and strain-hardening exponent to decrease with an increase in tempering temperature and holding time and the reduction in area and as elongation to increase with an increase in tempering temperature and holding time. After a temperature of 300°C toughness reduces due to retained inter-lath austenite and inter-lath carbide. Fracture analysis revealed that for all of the tempering temperatures, the fracture features are predominantly the ductile mode with a dimple structure. However at 300°C, the material failed in a brittle manner due to the occurrence of tempered martensite embrittlement.

S. Hossein Nedjad and co-workers [51] investigated the effect of aging microstructure on various mechanical properties of a Fe–Ni–Mn–Cr maraging alloy.

The material Fe–10Ni–7Mn–2Cr (wt.%) alloy exhibited a promising age hardenability but had intergranular brittle fracture. When chromium was added to the alloy it retained the austenite. However, improved alloy revealed a weak age hardenability. Age

hardening of the partially transformed Fe–10Ni–7Mn–6Cr (wt.%) alloy causes a remarkable increase in yield strength coupled with drastic loss of ductility.

CHAPTER III

THE TEST MATERIALS

The steels chosen for this study are AerMet ® 100, PremoMetTM 290, 300M and

TenaxTM 310. These materials were provided by Carpenter Technology,Inc based in

Reading Pennsylvania(Pa), USA. The composition of the four high strength steels chosen for this experimental study is provided in Table 3.1. Presence of carbon provides solid solution strengthening as well as hardenability through the formation and presence of alloy carbides. The alloy carbides serve to enhance the high temperature resistance and overall creep strength of a predominantly ferrite matrix. The presence of elements like chromium [Cr], molybdenum [Mo], Cobalt [Co] assists in the formation and presence of carbide particles, which contributes to enhancing the strength of the steel matrix.

However, the presence and distribution of carbide particles in the microstructure is detrimental to fracture toughness or impact resistance arising as a result of increasing the number of potential sites for the formation of fine microscopic cracks. Presence of nickel

[Ni] facilitates in lowering the transition temperature while concurrently enhancing toughness and stabilizing any austenite that is present in the material. The presence of molybdenum assists in refining the grain size, in addition to its role in forming molybdenum carbides and resultant influence in enhancing toughness [52].

The high strength steels designated by the manufacturer as AerMet® 100

(UNS K92580) and PremoMetTM 290 were produced through vacuum induction melting

(VIM) followed by vacuum arc re-melting (VAR) and subsequently cast as ingots. The high strength steels designated by the manufacturer as 300M (UNS K44220) and

TenaxTM 310 were produced by ARC/VAR processing [ARC being the initial air-melt operation to produce electrodes from subsequent VAR processing; VAR: Vacuum Arc

Re-melt]. AerMet® 100 is a high strength steel possessing high hardness and high strength coupled with exceptional ductility and toughness. This steel is used for designing components requiring high strength, high fracture toughness and exceptional resistance to stress corrosion cracking and fatigue [53]. The material has a nominal fracture toughness of 126 MPa √ (m) and tensile yield strength of 1724 MPa [53]. The steel designated as

300M has less strength and distinctly less fracture toughness than AerMet® 100 [53].

The PremoMetTM 290 steel and TenaxTM 310 steel are both free of cobalt. Further, these two steels have a combination of high strength and toughness in the quenched and tempered condition. These two steels attain an ultimate tensile strength of 2040 MPa coupled with a minimum fracture toughness of 77 MPa √ (m) and excellent fatigue life

[54]. The summary of the processing techniques on these specialty steels are explained in Table 3.2.

Table 3.1 Nominal chemical composition of the chosen materials

Material C Mn Si P S Cr Ni Mo Cu Co Al Ti Cb V N O

AerMet® 0.238 <0.01 0.03 0.002 0.0007 2.99 11.2 1.18 <0.01 13.4 0.003 0.011 0.003 - <0.001 <0.001 100

300M 0.43 0.75 1.7 0.005 <0.001 0.80 1.76 0.38 0.04 <0.01 0.01 <0.005 <0.01 0.08 - -

PremoMetT 0.404 0.79 1.50 0.003 0.0011 1.29 3.82 0.50 0.6 0.01 0.01 0.006 <0.002 0.30 0.002 <0.001 M 290 TenaxTM 0.401 0.63 2.03 0.004 <0.0005 1.26 3.76 1.01 0.53 0.01 0.01 0.003 <0.002 0.31 0.003 <0.001 310

Table 3.2 Primary and Secondary Processing history on the chosen High Strength Steels

Material Primary Processing Secondary Processing plus finishing AerMet® 100 (a) Melted by vacuum induction (i) The hot rolled bar (0.75 inch melting (VIM) and vacuum diameter) was annealed at 1250 F arc remelting (VAR) and cast (677oC) for 6 hours and subsequently to get ingots turned and ground to size. (b) Ingot was then hot-worked to (ii) The ground bar was tested and get 7.5 inch2 block. checked for defects using immersion (c) Block was then hot rolled to sonic inspection. get cylindrical bars having (iii)The optimum annealed hardness was 0.75 inch diameter. 40HRC (iv)Exhibits minimal size change during heat treatment.

PremoMetTM 290 (a) Melted by vacuum induction (i) The hot rolled bar (0.75 inch melting (VIM) and vacuum diameter) was annealed and arc remelting (VAR) and cast subsequently turned and ground to to get ingots size. (b) Ingot was then hot-worked to (ii) The ground bar was tested and get 7.5 inch2 block. checked for defects using immersion (c) Block was then hot rolled to sonic inspection. get cylindrical bars having 0.75 inch diameter.

TenaxTM 310 (a) Melted using vacuum arc (i) The hot rolled bar (0.75 inch remelting (VAR) and cast to diameter) was annealed and get ingots. subsequently turned and ground to (b) Ingot was then hot worked to Size. Get 7.5 inch2 block. (ii) The ground bar was tested and ( (c) Block was then hot rolled to checked for defects using immersion get cylindrical bars having Sonic inspection. 0.75 inch diameter.

300M (a) Melted using vacuum arc (i) The hot rolled bar (0.75 inch remelting (VAR) and cast to diameter) was austenitized at 870°C get ingots. and quenched at room temperature (b) Ingot was then hot worked to followed by double tempering at 315° get 7.5 inch2 block. C and subsequently turned and ( c (c) Block was then hot rolled to ground to size. get cylindrical bars having (ii) The ground bar was tested and 0.75 inch diameter checked for defects using immersion sonic inspection.

CHAPTER IV

TEST SAMPLE PREPARATION

4.1 Impact tests

Charpy-V Notch test specimens of the four steels were prepared in accordance with specifications outlined in ASTM E-23 [55](American Society for Testing Materials,

2006). The notched test specimens measured 55 mm in length, 10 mm in width and 10 mm in height. The notch was machined at the center which had 2 mm depth at an angle of 45 degrees. The key dimensions of the test specimens are shown in Figure 4.1.

4.2 Tensile tests.

Cylindrical test specimens, conforming to specifications outlined by American

Society for Testing Materials, ASTM E8-06 [56], were machined for the four chosen high strength steels. The threaded test specimens measured 120 mm in total length and

12.57 mm in diameter at the thread section. The gage section of the test specimen measured 25 mm in length and 6.25 mm in diameter. To minimize the effects of surface irregularities, the gage sections of the machined test specimens were at first mechanically ground on progressively finer grades of silicon carbide (SiC) impregnated emery paper and then finish polished, to a mirror-like finish, using an alumina-based polishing compound. The purpose of polishing was to remove any and all of the circumferential scratches and surface machining marks. The key dimensions of the test specimens are shown in Figure 4.2.

Figure 4.1: A schematic showing dimensions of the Charpy V-Notch test specimen.

Figure 4.2. A schematic of the cylindrical test specimen used for mechanical testing(Tensile).

CHAPTER V

EXPERIMENTAL TEST PROCEDURES

5.1 Microstructural Characterization

An initial characterization of the microstructure of the as-provided materials was done using a low magnification optical microscope. Samples of desired size were cut from the as-received material i.e. steels and were mounted in epoxy. The mounted samples were then polished using a series of silicon carbide impregnated emery paper

(240 grit,320 grit,400 grit and 600 grit) with water both as a lubricant and a coolant.

Subsequently, the four steel samples were mechanically polished using five-micron alumina solution and one-micron alumina solution. Fine polishing to a perfect mirror-like finish of the surface of all four steels was achieved using 0.1 micron alumina solution as the lubricant. The polished samples were subsequently etched using a reagent that is a solution mixture of 5-ml of nitric acid (HNO3) and 90 ml of Ethanol and this reagent is called as „Nital‟. The polished and etched surface of the four steel samples was observed in an optical microscope and photographed using standard bright field illumination technique.

5.2 Characterization of Microstructure using Atomic Force Microscopy (AFM).

A commercial digital instrument Atomic Force Microscope (AFM) [Model:

VEECO 3100] equipped with a Nanoscope IV controller was used to study the surface of the four candidate steels (AerMet® 100, PremoMetTM 290, 300M and TenaxTM 310). It

24 is to be noted that the surface of these high strength steels do not have the same uniform flatness at the microscopic level and approaching the nano-level. This is essentially ascribed to be due to heat treatment and related processing applied to each steel. Test sample of each high strength steel was rough polished using silicon carbide impregnated emery paper (from 240 grit to 600) Subsequently, the sample surface was finish polished using alumina-based polishing compound suspended in water to get a mirror-like finish.

The polished sample was etched using nital reagent (a solution mixture of nitric acid

(10ml)) in ethanol (90 ml). The atomic force microscope (AFM) provides a three- dimensional profile of the sample surface at the nanoscale, by measuring the forces between the probe and the polished sample surface at a short distance of 0.2-10 nm. The probe is supported on a flexible cantilever [57]. The tip of the AFM touches the polished surface of the metal gently and the small forces between the surface and the probe are recorded. The probe is connected at the end of a cantilever and the amount of force recorded by the probe depends on two key factors: (i) stiffness of the spring, and (ii) height of the spring [58]. Since polished surface of the steel is harder than the spring the cantilever tends to bend and resultant force on the tip is repulsive. Thus, the deflection is made constant and is used as a feedback loop while force between the probe and the sample surface remains constant. Finally, an image of the polished metal surface was obtained. The method used for measuring the force is known or referred to as the

“Contact Method” [57].

5.3 Hardness testing

Hardness of a material is a mechanical property defined as the resistance offered by the material to indentation i.e. permanent deformation and cracking.

5.3.1 Microhardness testing

The Vickers microhardness (Hv) of the four steel samples was measured using a

Wilson Tukon 2100B. The diamond indenter has a square-based pyramidal geometry with an included angle of 136 degrees. The machine makes an indent, or impression, on the sample surface whose diagonal size(s) was measured using a low magnification optical microscope. The indentation load used was 0.5 kgf. Use of lower loads with respect to high strength steels was not considered since it can cause problems due to load dependence of hardness coupled with measurement uncertainty arising from small indentation size [59]. The Vickers hardness number (Hv) is expressed as the ratio of applied load to surface area of indent. To minimize the uncertainties arising from the conjoint influence of measurement and cracking of the samples, an indentation load of

0.5 kgf for a dwell time of 15 seconds was chosen. At least four indents in each direction

(longitudinal and transverse) were made on the polished surfaces of the steel sample and the result is reported as the average value in units of kg/mm2. The hardness values are summarized in Table 5.1.

Table 5. 1 A compilation of the micro-hardness measurements made on the Specialty steels

Avg Trail Material Trail 2 Trail 3 Trail 4 Trail 5 Hardness 1 (kg/mm2) D1(μm) 39.62 39.62 39.62 38.7 39.39 AerMet® D2(μm) 39.6 39.83 39.6 38.66 39.36 100 Hv 590.94 587.46 590.94 619.76 597.95 597.41 (kg/mm2) D1(μm) 39.39 39.16 39.16 39.39 38.47

PremoMet D2(μm) 40.07 39.36 39.83 39.83 38.89 TM 290 Hv 587.42 601.49 594.37 590.9 619.72 598.78 (kg/mm2) D1(μm) 39.16 39.16 38.93 39.39 38.7

D2(μm) 39.6 39.83 38.66 39.36 38.89 300M Hv 597.91 594.37 616.06 597.95 616.02 604.462 (kg/mm2) D1(μm) 39.16 38.7 39.16 38.47 38.7

TM D2(μm) 39.13 39.13 39.36 38.89 38.89 Tenax 310 Hv 605.09 612.32 601.49 619.72 616.02 610.928 (kg/mm2)

5.3.2 Macrohardness testing.

The Rockwell hardness (RC) test is a static indentation test similar to the Brinell indentation test. It diﬀers from the brinell test in that it measures the permanent increase in indentation depth from the depth reached following an initial load of 98.1 N, upon the application of an additional load. Measurement is made after recovery, which takes place following the removal of the additional load. The Rockwell hardness number is a directly read (in units of 0.002 mm) from the dial gauge which is attached to the

Rockwell machine while an initial minor load is still imposed [60]. For the steels used in this study an indentation load of 150 Kgf , 120 degree diamond cone and a dwell time of

10 seconds were the variables used and the hardness value read on the “C‟ scale. The macro-hardness tests were also done on the polished surface of the steel specimens .

The Rockwell scales are divided into 100 divisions, each equivalent to 0.002 mm of recovered indentation. Since the scales are reversed, the number is higher if the material is harder.

Formula, which is used to find the hardness number in a Rockwell apparatus is

Rockwell C = (RC )= 130 –(depth of penetration (mm)/0.002) [60]

The results of the macro-hardness tests are summarized in Table 5.2.

Table 5. 2 A compilation of the macrohardness measurements made on the Specialty steels samples.

Trail Trail Trail Trail Trail Avg Material 1 2 3 4 5 Hardness

Rc 55 54 55 56 55 55.124 Hardness 301 299 301 314 301 303 AerMet® (ksi) 100 Hardness 211.8 210.4 211.8 221 211.8 213.2 (kg/mm2)

Rc 55 55 55 55 56 55.2 Hardness PremoMet 301 301 301 301 314 304 TM (ksi) 290 Hardness 2 211.8 211.8 211.8 211.8 221 214 (kg/mm )

Rc 55 55 56 55 56 55.4 Hardness 301 301 314 301 314 306 300M (ksi) Hardness 211.8 211.8 221 211.8 221 215.3 (kg/mm2)

Rc 56 56 55 56 56 55.6 Hardness TM 314 314 301 314 314 311 Tenax (ksi) 310 Hardness 221 221 211.8 221 221 219 (kg/mm2)

5.4 Impact Tests.

Charpy V-Notch specimens were prepared in conformance with specifications outlined in ASTM E-23 [55](Figure 1). Three replicate samples were tested at each of the six chosen test temperatures. In all, a total of 18 specimens were deformed under conditions of impact loading. Test samples were brought to the desired temperature by immersing in environments of:

(a) liquid nitrogen (T = -180°C)

(b) dry ice (T = -55°C)

(d) room temperature (T = 25°C)

(e) boiling water (T = 98°C)

(f) furnace (T = 170°C).

The samples were immersed in a specific environment for full 30 minutes prior to the initiation of testing. Specimens evaluated at the higher test temperature (T = 170°C) were placed in a Blue-M Moldatherm box furnace and allowed to soak at the temperature for well over 30 minutes prior to testing.

The Charpy V-Notch specimens were removed from the respective environments and quickly placed in the test fixture of the impact test machine (Model: Tinius-Olsen) with a capacity of 300 ft-lbs.This was followed by a quick release of the impact hammer, i.e., the pendulum. The total energy absorbed by the sample to failure was directly read from the graduated scale on the test machine.

5.5 Tensile Tests

Uniaxial tensile tests were conducted on an INSTRON-8500 Plus which is a closed-loop, fully automated servo-hydraulic mechanical test machine with a 100 kN load cell. The tests were performed at room temperature (298 K) and in the laboratory air environment having a relative humidity of 55 pct.. The steel test specimens were deformed at a constant strain rate of 0.0001/sec. An axial 12.5-mm gage-length clip-on extensometer was attached to the test specimen, using rubber bands, to provide a measure of the strain during uniaxial stretching The stress and strain measurements, parallel to the load line, were recorded on a PC-based data acquisition system (DAS).

5.6 Failure-Damage analysis:

Fracture surfaces of the fully deformed and failed samples under uniaxial loading and impact loading were examined comprehensively in a scanning electron microscope

(SEM). This is done to determine the fracture mode the material has undergone and is also characterizes the fine scale topography and sub grain features on the tensile fracture surface. This study helps in determining the microscopic mechanisms governing fracture during the tensile deformation and provides useful information related to the same.

The distinction between the macroscopic mode and microscopic fracture mechanisms is based on the magnification level at which the micrographs were taken.

The overall nature of failure is referred to as the macroscopic mode while the failure processes occurring at the local level is referred to as microscopic mechanisms, such as, microscopic void formation, microscopic void growth and eventual coalescence, and

nature, intensity (i.e.,number) and severity of the fine microscopic cracks and macroscopic cracks. The samples for observation in the scanning electron microscope

(SEM) were obtained from the deformed and failed specimens by sectioning parallel to the fracture surface [51,52]

CHAPTER VI

RESULT AND DISCUSSION

6.1 Initial Microstructure

The microstructure of any material is an important factor to be taken note of in any alloy. Properties like tensile properties, fracture toughness, fatigue resistance and resultant fracture behavior could be analyzed and it helps in giving an idea of how the material would be and what applications could it be used for. The optical micrographs of

AerMet® 100 and PremoMetTM 290 are shown in Figure 6.1 and Figure 6.2. The micrographs of 300M and TenaxTM 310 respectively are shown in Figure 6.3 and Figure

6.4 [53,54]. The micrographs are at three different magnifications and reveal the microstructure of the high strength steels to be a combination of martensite and ferrite. A higher carbon content in the steel resulted in a greater volume fraction of martensite. The presence and overall significance of martensite was more noticeable in large numbers and as lath morphology for both AerMet® 100 and PremoMetTM 290 that were produced by vacuum induction melting (VIM) and vacuum arc re-melting (VAR). The high strength steels 300M and TenaxTM 310 that were produced by vacuum arc re-melting (VAR) there was nothing significant by observable magnifications (Figure 6.3(a) and 6.4(a)) since it is very closely packed and the martensite was much finer and intermingled with pockets of randomly distributed ferrite-rich regions distributed randomly, i.e., carbon depleted regions in the microstructure. The presence of these features is governed by the

33 combination of composition and primary processing and does exert an influence on impact, fracture toughness and tensile properties of the candidate steels.

6.2 Initial Microstructure using the Atomic Force Microscope.

The atomic force microscope (AFM) was used to establish the surface abnormalities and size of the second-phase particles present in the material. The reasons for the presence of these abnormalities are the following:

(i) Alloying element (inclusions).

(ii) Heat treatment.

(iii) Secondary processing of the material.

(iv) The finishing process.

The images obtained from the atomic force microscope had a sectional size of 10 μm x

10 μm. The „contact method‟ was used to obtain the images. This method was chosen and used primarily because it is fast and good for rough sample surfaces. However, should the samples be soft then the „contact method‟ can break the sample. This is not applicable for the chosen steels since their hardness is noticeably greater than the cantilever of the microscope. The portion closest to the tip of the AFM is taken to be the origin and given a value 0 nm. The probe is then calibrated with specific reference to this position and the resultant low‟s and high‟s present on the surface is recorded. A colored legend was used to define the heights at various points of the chosen section of the material (i.e., high strength steel). The observed difference in height through the section of the chosen sample surface was only a few hundred nanometers.

The steels AerMet® 100 (Figure 6.5) and PremoMetTM 290 (Figure 6.6) were processed under similar conditions, namely, vacuum induction melting (VIM) followed

34 by vacuum arc re-melting (VAR). Thus, these two steels reveal a smaller difference in their height of surface undulations over the sample surface, with the maximum change being 200 nm. The high strength steels 300M (Figure 6.7) and TenaxTM 310 (Figure 6.8) were processed by vacuum arc re-melting (VAR) and had a rough surface revealed by the change in height across the section to be above 200 nm. This makes these two steels susceptible to heat treatment when compared with the two steels AerMet®100 and

PremoMetTM 290. Also, TenaxTM 310 was found to have a rougher surface than the other three high strength steels examined in this research study. While not much of information was obtained from this analysis using the AFM, it did help in understanding the specific influence of secondary processing on a material.

(a) (b )

100μm 20μm

(c)

10μm

Figure 6.1 Optical micrographs showing the key micro-constituents of AerMet® 100 at three different magnifications

(a) (b)

100μm 20μm

(c)

10μm

Figure 6.2 Optical micrographs showing the key micro-constituents of PremoMetTM 290 at three different magnifications

(a) (b )

100μm 20μm

(c)

10μm

Figure 6.3 Optical micrographs showing the key micro-constituents of 300M at three different magnifications

(a) (b )

100μm 20μm

(c)

10μm

Figure 6.4 Optical micrographs showing the key micro-constituents of Tenax TM 310 at three different magnifications

Figure 6.5 Atomic force microscope image of the etched surface of AerMet® 100 showing profile and nature of roughness on a 10 μm x 10 μm section.

Figure 6.6 Atomic force microscope image of the etched surface of PremoMetTM 290 showing profile of roughness on a 10 μm x 10 μm section.

Figure 6.7 Atomic force microscope image of the etched surface of 300M showing nature and profile of the roughness on a 10 μm x 10 μm section .

Figure 6.8 Atomic force microscope image of the etched surface of TenaxTM 310 showing nature and profile of roughness on a 10 μm x 10 μm section.

6.3 Hardness

Hardness is defined as the amount of resistance a material provides to permanent indentation. It is the easiest way to quantify certain mechanical properties of a material and it is not a global phenomenon of the material but provides the range over which the properties be within, thereby, this makes it simple, cost effective and non-destructive in knowing the mechanical properties of a material. The observed of intrinsic microstructural features in the material due to the heat treatment applied to the material.

6.3.1 Microhardness measurement

Microhardness testing can be defined as indentation hardness testing that involves forcing a diamond indenter of specific geometry onto the surface of the test material at varied loads depending on the material. The microhardness measurement is generally made from edge-to-edge across the sample which is mounted on epoxy. A few trials were made in order to gather a detailed set of readings over the total length of the sample and show the variation along the sample. The microhardness measurements serves to provide the net effect of strengthening arising from a combination of composition and applied to the material , and a weakening effect resulting from the presence of any processing- related artifacts. The presence of processing-related artifacts, such as, the fine microscopic pores and voids, and fine microscopic cracks, when intercepted by the pyramidal indenter can result in a net decrease in the value of measured microhardness of the sample. The microhardness values of the various steels are summarized in Table 5.1.

All the four chosen specialty steels materials had a high microhardness with an average value of 598 kg/mm2. TenaxTM 310 leading the list by a marginal amount. All the four steels were very relatively close to each other in terms of mircrohardness. This is

44 attributed to the manufacturing and heat treatment processes which was used for the steels and discussed in chapter II and chapter III.

6.3.2 Macrohardness measurement

The macrohardness values (Table 5.2), are summarized on the Rockwell C scale across the length of the machined steel specimens and ranges from 210 kg/mm2 to 220 kg/mm2 with an average value of 215 kg/mm2. Again Tenax TM 310 had the highest hardness with the Rockwell C hardness of 56, which is equal to 219 kg/mm2. The microhardness and macro-hardness values across the length of the test specimen of

AerMet®100 and PremoMet TM290 is shown in Figure 6.9 and Figure 6.10, these two figures reveals the microhardness values at an indentation site to be noticeably higher than the macrohardness value. A similar same effect is observed when we compare 300M and Tenax TM310 in Figure 6.11 and Figure 6.12. The observed lower value of the macrohardness of the steels is ascribed to the presence of a population of processing- related defects in the material. The average value of hardness of the four high strength steels were compared in the bar graph shown in Figure 6.13, which reveals the local hardness(micro-hardness), value of the samples to be nearly three times the value of the global hardness(macro-hardness), as expected.

Figure 6.9 A profile showing the variation of microhardness and macrohardness value across the length of AerMet® 100.

Figure 6.10 A profile showing the variation of microhardness and macrohardness values across the length of PremoMetTM 290 .

Figure 6.11 A profile showing the variation of microhardness and macrohardness values across the length of 300M .

Figure 6.12 A profile showing the variation of microhardness and macrohardness values across the length of TenaxTM 310 .

700

2 600 AerMet 100 500

400 PremoMet 290

300 300M 200 Tenax 310

100 Hardness in in kg/mmin in Hardness 0

Micro- Macro- Hardness Hardness

Figure 6.13 Bar graph depicting the average microhardness and macrohardness values of the four specialty steels .

6.4 IMPACT TOUGHNESS

The results of CVN impact tests on the four ultra-high strength steels reveal the conjoint influence of test temperatures and presence of notch on dynamic mechanical properties and fracture behavior. The steels showed an increase in absorbed energy with an increase in test temperature (Table 6.1). The observed increase confirms the validity of this test since ductility of a metal increases with an increase in temperature. A comparison of the tensile properties of each specialty steel with respective impact test results provides a direct correlation of relative performance. Steels (AerMet® 100 and

PremoMetTM 290) follow a similar trend but AerMet®100 had the highest capacity for energy absorption, as shown in Figure 6.14, AerMet®100 is consistent throughout the range of temperature. From -180°C to -50°C the values of energy absorbed (N-m) was almost the same but at the higher temperatures the values were different and AerMet®

100 was observed to be tougher than PremoMetTM 290. The change in energy absorption with an increase in test temperature is larger for AerMetTM 100 than for PremoMetTM

290. This observation reveals the AerMet® 100 to have better impact fracture toughness over the range of test temperatures. The rate of change of energy absorbed for the steels

AerMet®100 and PremoMetTM 290 is shown in Figures 10. The two other steels namely

300M and TenaxTM 310(Figure 6.15) which were processed similar manner follow the same trend and the amount of energy absorbed prior failure is almost the same through the range of temperatures. Steel 300M had lower energy absorbtion till 100°C than

TenaxTM 310 however, at furnace temperature the energy absorbed was higher for 300M and revealed a linear rate of change, which is shown in Figure 6.15. Four steels followed a similar trend, which is, during the early stages prior to stable crack propagation

51 culminating in failure. In the region of stable crack propagation, energy absorbed, or toughness of the material, increases linearly with an increase in test temperature. Such an increase is ascribed to be due to an increase in localized microplastic deformation ahead of the propagating crack and the resultant intrinsic microscopic mechanisms governing fracture. During the later stages the high energy absorbing capability of each of the high strength steels is ascribed to be due to fracture by ductile tearing. The only difference between the sample in the early stage and sample in the later stage is the amount and/or severity of tearing associated with ductile failure.

At a given test temperature there is noticeable difference in impact toughness of the steels (AerMet® 100 and PremoMetTM 290) both produced by vacuum induction melting and vacuum arc re-melting. While the difference in impact toughness between the steels 300M and TenaxTM 310 both produced by Vacuum arc re-melting was minimal except failure at the furnace temperature. An overall comparison of the four steels at all test temperatures is shown in Figure 6.16.

Figure 6.14 Variation of energy absorbed (N-m) as a function of test temperature for the two steels AerMet® 100 and PremoMetTM290

Figure 6.15 Variation of energy absorbed (N-m) as a function of test temperature for the two steels 300M and TenaxTM 310

-200° -55° 0° 25° 95° 170°

TEST TEMPERATURE(Celcius)

Figure 6.16 Bar graph comparing the energy absorbed with test temperature of the four chosen materials:AerMet® 100, 300M, PremoMetTM 290 and TenaxTM 310 .

Table 6. 1 The impact toughness properties of the four high strength steels in N-m

MATERIAL Liquid Dry Ice Ordinary Room Boiling Furnace Nitrogen (N-m) Ice Temp water (N-m) T= -55oC (N-m) (N-m) (N-m) (N-m) T=-180oC T= 0 oC T= 25 oC T= 95 oC T= 170oC AerMet®100 19 30 53 56 58 58

300 M 15 23 31 32 37 41 PremoMet TM 15 33 34 36 37 40 290

TenaxTM 310 15 27 35 35 37 38

6.5 DYNAMIC FRACTURE

(a) Impact Fracture of AerMet® 100

The fracture surface features of AerMet® 100 deformed at -180°C are shown in

Figure 19. A vast majority of the fracture surface is covered with a population of fine microscopic voids and dimples. These features are reminiscent of globally ductile behavior even at this low a test temperature. The overall morphology of the test sample was essentially smooth with lack of observable macroscopic cracks and other features reminiscent of globally brittle behavior (Figure 6.17a). High magnification observation of

(a) revealed a population of voids of varying size along with dimples (Figure 6.17b). At the next higher allowable magnification of the scanning electron microscope regions of the micrograph shown in (b) revealed clearly the morphology of the fine microscopic voids and their distribution on the impact fracture surface (Figure 6.17c). The macroscopic and fine microscopic voids have little time to grow and eventually coalesce to form one or more macroscopic cracks

(Figure 6.17d). The presence of few yet isolated populations of fine microscopic cracks provides evidence of locally brittle failure mechanisms.

The fracture surface features of AerMet® 100 deformed at -55°C are shown in

Figure 6.18. The overall morphology of the test sample was essentially smooth with lack of observable macroscopic cracks and other features reminiscent of globally brittle behavior (Figure 6.18a). High magnification observation of (a) revealed a random distribution of fine microscopic cracks and higher magnification of (6.18b) shows a random distribution of fine microscopic voids and a healthy population of ductile dimples

(Figure 6.18c). In Figure 6.18d we can see the shape and size of dimples. The features of

AerMet® 100 at -55°C was similar to liquid nitrogen and did show globally ductile behavior.

At room temperature the deformation was more ductile and the material absorbed more energy than it did at the lower temperatures thereby showing a noticeable change in the curve. Figure 6.19a shown the ruptured morphology of AerMet® 100 which reveals ductile behavior of the steel AerMet® 100. In Figure 6.19b which is a higher magnification of a region adjacent to the notch shows dimples and voids of varying size.

Macroscopic cracks surrounded by shallow dimples of varying sizes are shown in Figure

6.19c Fine microscopic cracks were observed in the vicinity of the notch and overload was observed(Figure 6.19d)

The fracture surfaces of the CVN sample dynamically deformed at the highest test temperature used in this study, i.e., 170° C, are shown in Figure 6.20. Fracture was visibly rough at the macroscopic level (Figure 6.20a) with an absence of features reminiscent of globally brittle fracture. Careful examination of the fracture surface at

56 higher magnifications in the region immediately adjacent to the notch revealed fine microscopic cracks intermingled with a rich population of dimples; features reminiscent of the brittle and ductile failure mechanisms operating at the „local‟ level. At gradually higher magnifications was revealed the size, morphology and distribution of the fine microscopic cracks (Figure 6.20c). The overload fracture surface revealed a noticeable population of fine microscopic voids intermingled with shallow elongated dimples

(Figure 6.20d). Elongation of the dimples results from the occurrence of ductile shear failure or due to the presence of localized Mode II deformation field at the microscopic level.

Voids and Dimples

(a) (b)

200μm 20μm

5μm 2μm

Microcracks

Figure 6.17 Scanning electron micrographs of Aermet® 100 deformed under impact following exposure in liquid nitrogen at – 180°C, showing: (a) Overall morphology of failure. (b) High magnification of (a) showing a healthy population of microscopic voids and dimples. (c) Morphology of the fine microscopic voids and their distribution (d) Isolated fine microscopic cracks reminiscent of locally brittle failure mechanisms.

Microcracks

(a) (b)

200μm 20μm

5μm 2μm

Voids and Dimples

Figure 6.18 Scanning electron micrographs of AerMet® 100 deformed under impact following exposure in dry ice at – 55°C, showing: (a) Overall morphology of failure. (b) High magnification of (a) showing a random distribution of fine microscopic cracks. (c) High magnification of (b) showing a random distribution of fine microscopic voids and a healthy population of ductile dimples. (d) The size and shape of the dimples on the overload surface.

(a) (b)

200μm 100μm

5μm 2μm

Macroscopic Crack Microscopic Crack

Figure 6.19 Scanning electron micrographs of Aermet® 100 deformed under impact in the room temperature environment at 25°C, showing:

(a) (b)

200μm 50μm

20μm 2μm

Microscopic Crack Microscopic voids

Figure 6.20 Scanning electron micrographs of Aermet® 100 deformed under impact following exposure in a furnace at 170°C

(b) Impact Fracture of PremoMetTM 290

The scanning electron micrographs of this high strength steel are shown in Figure

6.21 to Figure 6.24 and support the observation made from measurement of impact energy absorbed as a function of test temperature. The fracture surface features associated with the Charpy V-notch test specimens that was dynamically deformed at the lowest test temperature (-180 C) is shown in Figure 6.21. Careful observation revealed failure to have initiated at the root of the notch and propagated outward radially (Figure

6.21a) giving an overall appearance of ductile failure at the macroscopic level.

Observation of the region surrounding the root of the notch at the higher allowable magnifications revealed a population of fine microscopic voids and an absence of features reminiscent of locally brittle failure (Figure 6.21b). The morphology and shape of the microscopic voids revealed them to be of varying size and randomly distributed through the fracture surface (Figure 6.21c). Even at this low a test temperature this high strength steel revealed features reminiscent of locally ductile mechanisms. The region of

„tearing‟ prior to catastrophic failure revealed isolated microscopic cracks resulting from the coalescence of the existing fine microscopic voids (Figure 6.21d). The microscopic cracks were surrounded by a healthy population of dimples are reminiscent of the locally operating brittle and ductile failure mechanisms.

The fracture surface of the samples which failed at -55°C showed reminiscence of brittle failure. The overall morphology of the failure was brittle (Figure 6.22a). Higher magnification of the fracture surface at end near the notch revealed a random distribution of microscopic voids of varying size and shape. In Figure 6.22c, microscopic cracks thus

62 portraying a locally brittle failure are shown. The region away from the notched revealed a mixture of both macroscopic and fine microscopic voids mixed with shallow dimples

The fracture surface features of samples failed at room temperatures showed glimpses of ductility observation. In Figure 6.23a is shown the overall morphology of the fracture surface while, magnification observation of this microscopic region close to the notch revealed a healthy population of microscopic voids and isolated microscopical cracks. A magnification observation of Figure 6.23b reveals the size and distribution of the voids (Figure 6.23c). In the region away from the notch reveals a distribution of shallow dimples intermingled with fine microscopic voids of varying sizes and shape

(Figure 6.23d).

The fracture surface features observed for the CVN sample deformed at the higher test temperature, i.e., 170C, are shown in Figure 6.24. Overall morphology of fracture was no different than the test sample that was deformed at the lowest temperature (-

180°C) (Figure 6.24a). The region immediately adjacent to the root of the notch showed voids of varying size, shallow and healthy distribution of dimples intermingled with isolated microscopic cracks (Figure 6.24b). These features are reminiscent of the locally operating ductile and limited brittle failure mechanisms. At even higher magnification careful examination of selected regions of the fracture surface revealed a fine array of microscopic cracks that were intermingled with microscopic voids (Figure 26c). At region far away from the root of the notch and well into the „tearing‟ domain the predominant feature covering the fracture surface was dimples and they were near identical in both size and shape but overall shallow in appearance (Figure 6.24d). On an

63 average the energy absorbed during impact at this test temperature (180°C) is 2.7 times greater than the energy absorbed by the sample dynamically deformed at -170°C.

(a) (b)

200μm 20μm

5μm 2μm

Figure 6. 21 Scanning electron micrographs of PremoMetTM 290 deformed under impact following exposure in liquid nitrogen at – 180°C, showing: (a) Overall morphology of failure as damage radiates outward from the notch. (b) The region adjacent to the notch showing a healthy population of fine microscopic voids. (c) High magnification of (b) showing the morphology and shape of the microscopic voids. (d) An array of fine microscopic cracks on the overload fracture surface.

(a) (b)

200μ 20μm m

10μm 2μm

Microscopic Crack Macroscopic voids

Figure 6. 22 Scanning electron micrographs of PremoMetTM 290 deformed under impact following exposure in dry ice at – 55°C, showing: (a) Overall morphology of the fracture surface. (b) High magnification observation in the region immediately adjacent to the notch showing a random distribution of microscopic voids. (c) Microscopic crack reminiscent of locally brittle failure. (d) A mixture of both macroscopic and fine microscopic voids in the region away from the tip of the notch

Microscopic cracks and voids

(a) (b)

200μ 20μm m

5μm 2μm

Microscopic voids

Figure 6.23 Scanning electron micrographs of PremoMetTM 290 deformed under impact in the room temperature (25°C) laboratory air environment, showing (a) Overall morphology of failure (b) The region adjacent to the notch showing a healthy population of fine microscopic voids and isolated microscopic cracks. (c) High magnification of (b) showing the morphology, size and distribution of the microscopic voids. (d) High magnification observation of ( c) showing shallow nature and distribution of the dimples intermingled with fine microscopic voids.

(a) (b)

200μm 50μm

10μm 5μm

Microscopic Crack Dimples

Figure 6. 24 Scanning electron micrographs of PremoMetTM 290 deformed under impact following exposure in a furnace at 170°C, showing: (a) Overall morphology of the impact fracture surface. (b) The region adjacent to the notch showing features reminiscent of ductile and brittle failure mechanisms. (c) An array of fine microscopic cracks intermingled with microscopic voids. (d) The size, shape and shallow nature of the dimples in the region away from the notch tip.

The scanning electron micrographs of the test samples deformed from the cryogenic temperature, i.e., -180°C to furnace temperature of 170°C are shown from

Figure 6.25 to Figure 6.28. These micrographs support the observation of an increase in energy absorbed by the steel when the test temperature is increased. The fracture surface features of the CVN test sample dynamically deformed at -180° C are shown in Figure

6.25. Overall morphology of failure (Figure 6.25a) fails to reveal brittle behavior of the material, Careful observation near the notch revealed a number of voids of varying size that were distributed randomly throughout the region(Figure 6.25b). At higher magnification in the region away from the notch revealed an array of fine microscopic cracks intermingled with pockets of cleavage facets. The cleavage facets were essentially flat and near featureless (Figure 6.25c). The overload region revealed a number of dimples and voids of varying size and shape (Figure 6.25d).

The micrographs of the failed 300M steel samples at -55°C are shown in Figure

6.26. The sample revealed no significant difference from the previous one but showed a marginal improvement in the brittle behavior. The overall morphology of the failed surface which revealed brittle behavior (Figure 6.26a) but the energy absorbed was much higher than the steel sample deformed at -180°C. Higher magnification observation near the notch shows a population of voids and dimples of varying sizes and shapes (Figure

6.26b). The high magnification of Figure 6.26b reveals the morphology and shape of microscopic void and dimples. In the region away from the notch was observed, shallow dimples of non-uniform sizes and shapes as shown in Figure 6.26d.

At room temperature the sample revealed ductility. The overall morphology of the failed sample is shown in Figure 6.27a which the crack radiates away from the notch . At higher magnification near the notch revealed a region reminiscent of ductile failure(Figure 6.27b). The void coalescence to form microcracks (Figure 6.27c) this region is surrounded by voids of varies sizes. A series of microcracks and shallow dimples were found covering in the region of overload (Figure 6.27d).

The fracture surface features of the CVN test sample dynamically deformed at the highest test temperature used in this study, i.e., 170oC are shown in Figure 6.28.

Macroscopically failure radiated outward from the root of the notch similar to the observation of the sample deformed at the lowest test temperature, i.e., -170°C (Figure

6.28a). High magnification observation in the region immediately adjacent to the notch revealed a healthy population of shallow dimples intermingled with voids of varying size, features reminiscent of locally ductile failure mechanisms (Figure 6.28b). The fine microscopic voids had limited time to grow and coalesce to form very fine microscopic cracks surrounded by a large population of essentially ductile dimples (Figure 6.28c).

Remnants of fine microscopic cracks intermingled with a random distribution of shallow dimples were evident in the region of „tearing‟ or overload (Figure 6.28d). These features are indicative of predominantly ductile with trace amounts of brittle failure mechanisms operative at the „local‟ level [53].

Microscopic voids

(a) (b)

200μm 20μm

5μm 2μm

Microscopic Crack Macroscopic voids and dimples

Figure 6.25 Scanning electron micrographs of 300M deformed under impact following exposure in liquid nitrogen at – 180°C, showing: (a) Overall morphology of failure. (b) High magnification observation in the region immediately adjacent to the notch showing voids of varying size and their random distribution. (c) High magnification observation of the impact fracture surface away from the notch. (d) A healthy population of shallow dimples of varying size and shape.

(a) (b)

200μm 20μm

5μm 1μm

Microscopic voids

Figure 6.26 Scanning electron micrographs of 300M deformed under impact following exposure in dry ice at – 55°C, showing: (a) The overall morphology of the impact fracture surface. (b) High magnification observation in the region adjacent to the notch showing a population of voids of varying size and healthy distribution of dimples. (c) High magnification of (b) showing the morphology and shape of microscopic void and dimples. (d) The shallow nature, and non-uniform size and shape of the dimples in the region away from the notch.

Macroscopic and microscopic voids

(a) (b)

200μm 10μm

2μm 1μm

Macroscopic voids and dimples Microscopic Crack

Figure 6.27 Scanning electron micrographs of 300M deformed under impact in the room temperature (25°C) laboratory air environment, showing: (a) The overall morphology of failure as it radiates away from the notch. (b) High magnification observation in the notch region showing a random distribution of macroscopic and fine microscopic voids. (c) The distribution and shape of shallow dimples and fine microscopic voids. (d) Isolated microscopic crack in a predominantly ductile region.

(a) (b)

200μm 20μm

10μm 2μm

Microscopic Crack

Figure 6.28 Scanning electron micrographs of 300M deformed under impact following exposure in a furnace at 170°C, showing (a) Overall morphology of failure as it radiates outward from the root of the notch. (b) High magnification observation in the region adjacent to the notch reminiscent of locally ductile failure mechanisms. ( c) Microvoid coalescence to form fine microscopic crack surrounded by dimples (d) Microcracks and random distribution of shallow dimples in the region of overload.

(d) Impact Fracture of TenaxTM 310

The fracture surface features of the CVN test sample that was dynamically deformed at the lowest test temperature of -180°C are shown in Figure 6.29. Overall morphology of fracture was essentially smooth and radiating outward from the root of the notch (Figure 6.29a). High magnification observation in this region, immediately adjacent to the notch tip, revealed a large population of dimples intermingled with both macroscopic and fine microscopic voids, features indicative or suggestive of „locally‟ operating ductile failure mechanisms (Figure 6.29b). Continued high magnification observation of this region revealed a sizeable population of microscopic voids, dimples of varying size and shape, and an isolated distribution of microscopic cracks (Figure 6.29c).

The dimples were shallow in nature. The limited growth and eventual coalescence of the closely spaced microscopic voids resulted in the formation of isolated microscopic cracks; few and far in-between (Figure 6.29d).

At -55°C the fracture surface showed a little ductile features even though it was brittle failure overall. The overall morphology of the failed sample is shown in Figure

6.30a and at higher magnification of the region next to the notch revealed voids of varying size and isolated microscopic cracks (Figure 6.30b). The growth of microscopic cracks to form macroscopic cracks and is surrounded by shallow dimples (Figure 6.30c).

The shape and size distribution of the dimples is shown in Figure 6.30d.

At room temperature the material was nearly completely ductile and reminiscence of ductile failure was observed throughout. The crack radiated outwards from the notch

(Figure 6.31a). At higher magnifications the ductile features could be observed distributed through the fracture surface of the material (Figure 6.31b). The size and shape

75 of microscopic voids are shown in Figure 6.31c, were which is high magnification observation of Figure 6.31b. The macroscopic cracks surrounded by dimples and voids

(Figure 6.31d).

The fracture surface features of the CVN test sample that was deformed under conditions of impact at the highest test temperature (180°C) are shown in Figure 6.32.

Overall morphology of the fracture surface was essentially flat and quite similar to the surface of the test sample deformed at the lowest test temperature (-170°C) (Figure

6.32a). The region immediately adjacent to the notch when viewed at gradually increasing magnifications of scanning microscope revealed a sizeable population of microscopic cracks, microscopic voids of varying shape intermingled with a sizeable population of shallow dimples (Figure 6.32b). The presence of these features is indicative of the locally operating ductile and brittle failure mechanisms. The profile or shape of a typical microscopic crack observed in this region is shown in Figure 6.32c. In the region of ductile „tearing; microscopic cracking was observed to be the dominant fracture feature and covered a sizeable portion of the fracture surface in this region (Figure

6.32d). The presence of these fine microscopic cracks is clearly indicative of locally brittle failure mechanisms. The transgranular region immediately adjacent to the macroscopic cracks was either flat or covered with pockets of shallow dimples. The impact energy absorbed by this test sample is 2.4 times greater than the energy absorbed by the CVN test sample deformed at the lowest test temperature (-180°C).

(a) (b)

200μ 10μm m

2μm 1μm

Microscopic Crack

Figure 6.29 Scanning electron micrographs of TenaxTM 310 deformed under impact following exposure in liquid nitrogen at – 180°C, showing: (a) Overall morphology of the fracture surface. (b) High magnification in the region adjacent to the notch tip. (c) High magnification of (b) showing a healthy population of microscopic voids, dimples and microscopic cracks. (d) Shallow nature of the dimples and coalescence of microscopic voids to form microscopic cracks.

Macroscopic Crack

(a) (b)

200μ 20μm m

10μm 2μm

Microscopic Crack Figure 6.30 Scanning electron micrographs of TenaxTM 310 deformed under impact following exposure in dry ice at – 55°C, showing: (a) Overall morphology of the fracture surface (b) High magnification in the region adjacent to the notch showing voids of varying size and isolated microscopic cracks. (c) High magnification of (b) showing growth and eventual coalescence of microscopic cracks to form macroscopic crack surrounded by a population of dimples. (d) The size, shape and distribution of dimples.

(a) (b)

200μm 10μm

5μm 2μm

Macroscopic Crack Figure 6. 31 Scanning electron micrographs of TenaxTM 310 deformed under impact in the room temperature (25°C) laboratory air environment, showing: (a) Overall morphology of the fracture as it radiates outward from the notch tip. (b) High magnification showing the features adjacent to the notch. (c) High magnification of (b) showing the size and distribution of the microscopic voids. (d) Macroscopic cracks surrounded by population of fine microscopic voids and dimples.

(a) (b)

200μm 20μm

5μm 4μm

Macroscopic Crack

Figure 6.32 Scanning electron micrographs of TenaxTM 310 deformed under impact following exposure in a furnace at 170°C, showing: (a) Overall morphology of the impact fracture surface. (b) The region immediately adjacent to the notch showing a population of microscopic cracks, microvoids and shallow dimples. (c) High magnification of (b) showing profile of a typical microscopic cracks. (d) Distinct macroscopic crack in the region away from the notch tip.

6.6 Tensile Properties

The tensile properties of the chosen materials, at the ambient temperature (25°C) are summarized in Table 6.2. Results reported are the average values based on duplicate tests.

(i) The elastic modulus of the chosen steels were identical and was between the range of

190 GPa and 198 GPa (Figure 6.33).

(ii) The yield strength of TenaxTM310 is 1791 MPa, which is the highest amongst the

chosen steels and AerMet®100 is close by with 1760MPa, which is roughly 2% less.

steel 300M had a yield strength of 1679MPa which is 12% less than TenaxTM 310 and

PremoMetTM 290(1597 MPa) had the lowest yield strength among all the chosen four

steels (Figure 6.34).

(iii)The ultimate tensile strength (UTS) of the four chosen steels was only marginally

different and ranged from 1993 MPa for AerMet® 100 to 2166 for

TenaxTM 310 (Figure 6.34).

(iv) The ductility quantified by elongation over 1 inch (25 mm) gage length was

highest for AerMet® 100 (28 %) and lowest for Tenax TM 310 (10 %) of the four

chosen high strength steels. The observed decrease in ductility is commensurate with

the carbon content of the steel. A higher carbon content results in the formation and

presence of more carbide particles in the microstructure. These particles are

intrinsically brittle and easily conducive for early cracking and failure during

uniaxial tensile deformation.

(v) The reduction in test specimen cross-section area, a direct measure of ductility, was

highest for AerMet® 100 (65 percent) and lowest for 300M (35 percent) A bar chart

comparing the ductility of the chosen steels is shown in Figure 6.34.

(vi) The tensile ductility, defined as ln (Ao /Af), was highest for AerMet® 100 (105 pct)

and lowest (45 pct.) for 300 M.

The engineering stress versus engineering strain curves for the four steels AerMet® 100,

300M, PremoMetTM 290 and TenaxTM 310 are compared in Figure 6.35. The test samples chosen from the four steels exhibited nearly identical elastic behavior and near similar strain-hardening characteristics. Immediately following the onset of necking, i.e. ultimate tensile strength, all the steel samples exhibited large plastic strains prior to failure. The

Strain hardening exponent values for these steels were found by drawing the true stress and true strain curves (Figure 6.36 to Figure 6.39) and the value of K was found which is the true stress at 1% true strain. This value is used in finite element analysis (FEA) for these steels. The results obtained from the true stress-true strain graphs are tabulated and presented in Table 6.3.

Elastic Modulus (GPa)

AerMet®100 300M PremoMetTM290 TenaxTM310

Figure 6.33 Bar graph depicting the elastic modulus of the four candidate high strength steels.

Strength (MPa)

AerMet®100 300M PremoMetTM290 TenaxTM310

Figure 6.34 Bar graph comparing the Yield Strength (MPa) and Ultimate TensileStrength (MPa) of the four chosen and studied high strength steels

Table 6. 2 A compilation of the room temperature tensile properties of the test Materials (four high strength steels) for the independent tests on the samples.

Elastic Modulus Yield Strength UTS Tensile Elongation Reduction in Ductility GL=1” Area ln(Ao/Ar) Specimen ksi GPa ksi MPa ksi MPa (%) (%) (%)

AerMet 100 27539.64 189.88 255.25 1759.79 289.05 1992.92 27.6 66.32 104.85

300M 28418.85 195.94 243.48 1678.72 297.21 2049.17 18.88 35.39 44.63

PremoMetTM 28698.53 197.87 231.57 1596.62 292.17 2017.46 20.99 45.24 60.22 290 TenaxTM 310 28059.60 193.46 243.09 1676.03 301.21 2076.80 N/A N/A N/A

Figure 6.35 The engineering stress versus engineering strain curve for the four high strength steels deformed in tension at room temperature (T = 25°C).

Table 6. 3 A summary of the monotonic mechanical parameters of the candidate high strength steels.

Material n k

AerMet® 100 0.157 1963.67

300M 0.109 1970.00

PremoMet™ 290 0.016 1894.54

Tenax™ 310 0.124 2031.93

10000 T = 25oC AerMet® 100 1000 n=0.16

100

10 True Stress(MPa) True

1 0.1 1 True Strain(%)

Figure 6.36 Variation of true stress (MPa) with true strain (%) for AerMet® 100 deformed in tension at room temperature (T=25°C). .

10000 T = 25oC PremoMet™ 290 1000 n=0.017

100 True Stress(MPa) True 10

1 0.1 1 True Strain(%)

Figure 6. 37 Variation of true stress (MPa) with true strain (%) for PremoMet™ 290 deformed in tension at room temperature (T = 25oC).

10000 T = 25oC 300M

1000 n=0.11 100

True Stress(MPa) True 10

1 0.1 1 True Strain(%)

Figure 6.38 Variation of true stress (MPa) with true strain (%) for 300M deformed in tension at room temperature (T = 25oC).

10000 T = 25oC Tenax™ 310

1000 n=0.11

100

True Stress(MPa) True 10

0.1 True Strain(%) 1

Figure 6.39 Variation of true stress (MPa) with true strain (%) for Tenax™ 310 deformed in tension at room temperature (T = 25oC).

6.7 Tensile Fracture Behavior

The tensile fracture surfaces of the four chosen steels were examined in a scanning electron microscope (SEM) to provide useful information relating to the specific role of intrinsic microstructural features and microstructural effects on strength, ductility and fracture properties of the candidate steel.

(a) Tensile Fracture of AerMet® 100:

Scanning electron micrographs of the tensile fracture surface revealed fracture to occur by the characteristic “cup and cone” type of separation which is a basic sign for ductile failure (Figure 6.40a). On higher magnification observation of the fracture surface reveals noticeable population of voids of varying size inter dispersed with a rich array of ductile dimples on the fracture surface (Figure 6.40b). The region of overload revealed a heterogeneous distribution of fine microscopic voids intermingled with fine microscopic cracks (Figure 6.40c). These features are indicative of the locally operating ductile and brittle fracture mechanisms. During far-field loading the microscopic voids appeared to have undergone limited growth and with time they eventually coalesce to form one or more microscopic cracks. Coalescence occurred from the progressive linkage of the larger macroscopic voids with the finer size and more closely spaced microscopic voids, through the formation of void sheets. The halves of these voids, both macroscopic and fine microscopic are the shallow dimples of varying size visible in large number on the tensile fracture surface (Figure 6.40d).

(b) Tensile Fracture of PremoMetTM 290:

Scanning electron microscopy observations of the tensile fracture surface revealed the occurrence of pronounced necking prior to failure and commensurate with the ductility of this specific steel (Figure 6.41a). Progressive higher allowable magnifications of the electron microscope revealed a large portion of the tensile fracture surface to be essentially transgranular. The transgranular regions were flat and near featureless (Figure

6.41b). Higher magnifications revealed the transgranular region and the regions immediately adjacent to the transgranular region to contain a population of fine microscopic cracks intermingled with fine microscopic voids reminiscent of “locally‟ operating brittle and ductile failure mechanisms (Figure 6.41c). The region of overload was covered with a sizeable population of fine microscopic voids of varying size and shape. This region also revealed an adequate number of cracked second-phase particles intermingled between the growth and eventual coalescence of the finer microscopic voids to form a microscopic crack (Figure 6.41d). Half of a void is the dimple that is visible on the tensile fracture surface. The presence of voids of varying size results in shallow dimples of varying size and shape, and present in large numbers on the overload fracture surface providing reminiscence of locally operating ductile failure mechanism.

Scanning electron microscopy observation, conducted with care, caution and comprehensively, of the tensile fracture surface revealed fracture to have occurred by a

“cup-and-cone” type separation (Figure 6.42a). Higher permissible magnifications of the scanning electron microscope revealed the tensile fracture surface to be covered with a population of heterogeneously distributed microscopic voids intermingled with a sizeable number of dimples (Figure 6.42b). The region of overload revealed cracking along the grain boundaries and at the grain boundary triple junctions, The observed features are commensurate with the noticeably lower ductility of this high strength steel (Figure

6.42c). At the higher allowable magnification of the electron microscope the grains were found to be covered with pockets of shallow dimples and fewer, yet noticeable population of fine microscopic voids. These features are reminiscent of locally ductile failure mechanisms (Figure 6.42d).

(d) Tensile Fracture of TenaxTM 310

Scanning electron micrographs of the tensile fracture surface revealed overall morphology to be essentially transgranular (Figure 6.43a). At the higher magnifications the transgranular region revealed a random distribution of very fine microscopic cracks

(Figure 6.43b). The region of tearing or overload failure revealed a mixture of voids of varying size intermingled with pockets of non-uniform size and shaped dimples (Figure

6.43c). The features observed at the higher magnifications of the scanning electron microscope, that is, fine microscopic cracks, microscopic voids and dimples are suggestive of the locally operating ductile and brittle failure mechanisms. The extent and

91 severity of the observed microscopic cracks was noticeably higher in view of the lower ductility of this high strength steel.

6.8 Mechanism Governing Tensile Fracture

During far-field loading in tension the presence of dislocation pile-ups both at the grain boundaries and the coarse second-phase particles present and distributed through the matrix does assist in the early initiation of fine microscopic voids at the second-phase particles in the microstructure. This is particularly favored to occur when the

“local”strain caused by dislocation pile up at the matrix-second phase particle interface reaches a critical value. During far-field tensile deformation a few of the second phase particles are favored to fracture on account of their intrinsic brittleness. This is aided by a

“local” reduction in the strain energy required for cracking, or separation, from the metal matrix. Since crack extension under quasi-static loading is favored to occur at high local stress intensities comparable to the fracture toughness of the material, the presence of a population of macroscopic voids and fine microscopic voids degrades the actual strain-to- failure associated with ductile fracture.

At the microscopic level, the formation and presence of a sizeable population of fine microscopic voids essentially transforms the deforming high strength steel into a composite material at the microscopic level comprising two populations of particles (i) the grains in the matrix, and (ii) voids (a void being considered as a particle having zero stiffness). Since the voids are intrinsically softer than the hardened grains in the matrix, the local strain is significantly elevated at and around the region of a microscopic void causing or enabling conditions that facilitate an increase in their volume fraction.

(a (b) )

200μm 50μm

(c (d) )

10μm 1.25μm

Figure 6. 40 Scanning electron micrographs showing tensile fracture surface of AerMet®100 deformed at room temperature (25° C), showing: (a) Overall morphology of failure. (b) High magnification of (a) showing a healthy population of voids of varying size inter-dispersed with a population of dimples (c) High magnification observation of the region of overload showing microscopic voids and microscopic cracks. (d) The shallow nature of dimples and microvoid coalescence to form microscopic cracks.

(a) (b)

200μm 50μm

10μm 2.5μm

Figure 6. 41 Scanning electron micrographs of the tensile fracture surface of PremoMetTM290 deformed at room temperature (T = 25°C), showing:

(a) (b)

200μm 50μm

5μm 2μm

Figure 6. 42 Scanning electron micrographs of the tensile fracture surface of 300M deformed at room temperature (T = 25°C), showing: (a) Overall morphology of failure (b) High magnification of (a) showing a healthy population of microscopic voids intermingled with dimples. (c) High magnification observation in the region of overload showing cracking at grain boundary triple junction. (d) Grains covered with shallow dimples and microscopic voids, reminiscent of locally ductile failure mechanisms.

(a) (b)

4μm 2μm

(c)

1μm

Figure 6.43 Scanning electron micrographs of the tensile fracture surface of Tenax TM 310 deformed at room temperature (T = 25°C), showing:

CHAPTER VII

CONCLUSIONS

7.1 Impact Toughness.

In this detailed experimental study the influence of test temperature on impact toughness property and fracture behavior of four high strength steels having varying chemical composition and processing history was investigated and the following are key observations.

1. Microstructure is typical of high strength steels and reveals a combination of carbon-

rich region and a carbon-depleted region. A higher carbon and alloy content in the

steel resulted in a greater volume fraction of martensite in the carbon-rich region.

The presence and overall morphology and significance of martensite was noticeable

and present in large numbers in the form of “lath” in the steels AerMet® 100 and

PremoMetTM 290. The martensite was much finer and intermingled with random

pockets of ferrite-rich region, i.e. the carbon depleted region in the high strength

steels 300M and TenaxTM 310.

2. Charpy impact tests on each of the chosen high strength steel were conducted at

temperatures ranging from -180°C to +170°C. Each of the chosen high strength

steels revealed an increase in energy absorbed with test temperature. At a given test

temperature the difference in energy absorbed between PremoMetTM 290 and

AerMet® 100 was noticeably different and dictated by chemical composition. Also,

minimum to no difference in energy absorbed between 300M and TenaxTM 310

processed by vacuum arc re-melting.

3. For a given high strength steel the macroscopic fracture mode was flat at

all of the test temperatures the surfaces were examined in the scanning electron

microscope. At progressively higher magnification the fracture surface revealed a

sizeable population of dimples intermingled with fine microscopic voids of varying

shape and isolated microscopic cracks, features reminiscent of predominantly locally

ductile failure mechanisms with trace amounts of brittle failure mechanism.

4. Over the entire range of test temperatures examined the overall fracture surface

morphology and intrinsic microscopic features observed on the fracture surface was

found to be nearly identical for the four high strength steels.

7.2 Tensile Deformation and Fracture Analysis.

Based on the detailed study aimed at understanding composition and processing influences on microstructural development, tensile properties and fracture behavior of the four high strength steels, the following are the key findings:

1. The elastic modulus of the four chosen steels was identical and in the range between

190 GPa and 198 GPa. The yield strength of the four chosen steels was marginally

different. The ultimate tensile strength (UTS) of the four chosen steels was only

marginally different and ranged from 1993 MPa for AerMet® 100 to 2166 for

TenaxTM 310.

2. The ductility quantified by elongation over 0.5 inch (12.7 mm) gage length

was highest for AerMet® 100 (28 percent) and lowest for TenaxTM 310 (10 percent)

of the four chosen high strength steels. The observed decrease in ductility is

commensurate with the carbon content of the steel. A higher carbon content results in

the formation and presence of more carbide particles in the microstructure.

3. It was inferred that during tensile deformation the presence of dislocation pile-ups

both at the grain boundaries and the coarse second-phase particles present and

distributed through the matrix does assist in the early initiation of fine microscopic

voids at the second-phase particles in the microstructure. This occurs when the

“local” strain caused by dislocation pile up at the matrix-second phase particle

interface reaches a critical value.

4. From the true stress-true strain graphs, it was found that, AerMet®100 to have the

highest strain hardening exponent compared to the other three steels this was

followed by TenaxTM 310 with PremoMetTM 290 having longest strain hardening

exponent of the four steels.

5. The presence of a population of macroscopic voids and fine microscopic voids

degrades the actual strain-to-failure associated with ductile fracture.

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32. C.M Wayman, Introduction to the crystallography of martensite transformation ,MacMillan,New York,1964

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46 ASM Handbook, Vol. 8, 1990, pg. 124-142.

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103

APPENDIX A

PROCEDURE FOR PERFORMING THE TENSION TEST ON THE INSTRON-8500

SERVO HYDRAULIC TESTING MACHINE

Step 1: POWER.

1. The power switch on the surge bar should be on.

2. Water valve should be turned on.

3. Switch on the CPU and the screen.

4. Switch the actuator.

5. Press <1> button and wait for 5 seconds.

6. Then press <2>

7. Check for errors or warnings in the console on the screen.

Step 2: Calibrate the load.

1. Right click on the load module on the console, and click on “Calibrate”.

2. Click “Next” four times until the screen shows “Start” and click “Start” for the

calibration to complete.121

3. The calibrate signal in the load area will blink continuously and stop, upon

completion of CALIBRATION.

4. Uncheck the ticked boxes for overwriting the calibration settings and click

“Finish”

105

Step 3: Specimen Loading.

1. Right click and click

2. Fix the specimen into the grips, this can be done by fixing the upper end of the

specimen to the grips and then moving the actuator up until there is minimal gap and

then fix the lower end of the sample.

Step 4: Recalibrate the load with specimen attached.

1. This is done because, while fixing the specimen some load was exerted on the

actuators. To negate that load-effect, we need to re-calibrate the load.

2. Turn load protector

3. Right click on the load module on the console, and click on “Calibrate”.

4. Click “Next” four times until the screen shows “Start” and click “Start” for the

calibration to complete.

5. The calibrate signal in the load area will blink continuously and stop, upon

completion of CALIBRATION.

6. Uncheck the ticked boxes for overwriting the calibration settings and click

“Finish”.

Step 5: Calibration of the strain gage.

1. Attach the strain gage plug to the strain socket.

2. Right click on the load module on the console, and click on “Calibrate”.

3. Click “Next” four times until the screen shows “Start” and click “Start” for the

calibration to complete.

4. The calibrate signal in the load area will blink continuously and stop, upon

completion of CALIBRATION.

106

5. The calibration of the strain gage is done WITHOUT fixing the strain gage to the

specimen.

Step 6: Fixing the strain gage to the specimen.

1. Once the specimen is firmly in the grips and the strain gage has been calibrated,

the strain gage must be fixed onto the specimen at the center of the gage length of

the specimen (the region where failure is predicted).

2. Ensure that the PIN remains in the strain gage until both the arms of the strain

gage are fixed onto the specimen with the aid of rubber bands.

3. Right Click on the load module to turn the

4. Once the arms of the strain gage are fixed onto the specimen with the help of

rubber bands, remove the pin.

5. It is advisable that the pin be removed gently, so as not to disturb the strain

calibration.

6. The calibration can be done again to ensure the reading of strain on the console is

close to zero.

Step 7: Choosing the program to run the tension test.

1. Choose Bluehill 2.0 on the desktop.

2. Create a Test method: choose a tension test method present in the list as per the

sample and its geometry.

3. FileOpenTest methods: (Manigandan/Steels Tension test).

4. Click OK.

5. Check for the parameters of the tensile test method: Check all the specimen

dimensions, test parameters and data storage parameters.

107

6. Set Strain control rate = 0.6%/min (This will depend on the material being tested)

7. Once the settings are set up, return to the HOME SCREEN.

8. Select the TEST icon.

9. Enter a file name for the test. (Example: AerMet100_3)

10. Click OK and click next until the “Start Test” button is enabled.

11. Click on “Start Test” button. The test will begin and the graph on the computer

screen will show the variation of stress with strain (if chosen).

12. The test will stop when the sample fails.

13. Remove one of the two rubber bands; insert the PIN back into the strain gage and

remove the other rubber band, and place the strain gage safely in the box.

14. Right Click on the load module to turn the

15. Remove the specimen from the grips.

16. The computer is programmed to give the print out of the data and the graph for

reference.

17. Click on “Next” until “Finish Test” button is enabled. Click “Finish Test”.

18. Ensure that you see a green progress bar indicating that the raw data collected

during the test is saved to the specified location.

19. Press < low> on the actuator controller for 5 seconds and then switch .

20. Turn off the water supply and switch the computer CPU and the monitor.

108

APPENDIX B

PROCEDURE FOR PERFORMING LOOP SHAPING FOR A GIVEN STRESS RATIO

AND MATERIAL ON THE INSTRON-8500 SERVO HYDRAULIC TESTING

MACHINE

Step 1: Switch on the POWER.

1. Switch on the power switch on the power surge bar.

2. Turn on the water flow for the hydraulic plant.

3. Switch on the computer and the CPU.

4. Switch the actuator.

5. Press button for 5 seconds.

6. Then press

7. Check for the status in the message console on the monitor to make sure that it

does not show any errors or warnings.

Step 2: Calibrate the load.

1. Right click on the load module on the console, and click on “Calibrate”.

2. Click “Next” four times until the screen shows “Start” and click “Start” for the

calibration to complete.

3. The calibrate signal in the load area will blink continuously and stop, upon

109

completion of CALIBRATION.

4. Uncheck the ticked boxes for overwriting the calibration settings and click“Finish”.

Step 3: Specimen Loading.

1. Press

2. Fix the specimen into the grips, this can be done by fixing the lower end of the

specimen first and then moving the actuator so that the specimen is in place

and the then the upper end of the specimen can be fixed firmly into the grips.

Step 4: Recalibrate load with specimen attached.

1. This is done because, while fixing the specimen some load was exerted on the

actuators. To negate that load-effect, we need to re-calibrate the load.

2. Turn load protector

3. Right click on the load module on the console, and click on “Calibrate”.

4. Click “Next” four times until the screen shows “Start” and click “Start” for the

calibration to complete.

5. The calibrate signal in the load area will blink continuously and stop, upon

completion of CALIBRATION.

6. Uncheck the ticked boxes for overwriting the calibration settings and click

“Finish”.

Step 5: Loop shaping with a test specimen.

1. Right click on the load module on the console, and click on “General” Option.

2. Click on the “PID” tab and note down the values for the fields proportional,

Integral, derivative and Lag.

3. Click on the “Properties” tab and modify the value of the field, “Tuning

110

disturbance amplitude” to a small value say 0.5 kN.

4. Click on “Auto-tune now” and wait for the process to complete.

5. Now, Click on the “PID” tab and note down the values for the fields‟

proportional, Integral, derivative and Lag. Repeat steps 3 and 4 with gradually

incrementing the Tuning disturbance amplitude value until the values of the above

field remains more or less constant.

Step 7: Removal of specimen after Loop tuning.

1. Once the loop tuning is complete, enable specimen protect and carefully remove

the specimen.

2. Press < low> on the actuator controller for 5 seconds and then switch .

3. Turn off the water supply and switch the computer CPU and the monitor

111