processing-microstructure-microtexture-property

PROCESSING-MICROSTRUCTURE-MICROTEXTURE-PROPERTY CORRELATION OF DUPLEX STAINLESS STEELS

Arka Mandal

PROCESSING-MICROSTRUCTURE-MICROTEXTURE-PROPERTY CORRELATION OF DUPLEX STAINLESS STEELS

Thesis submitted to the Indian Institute of Technology, Kharagpur For award of the degree of

Master of Science by

Arka Mandal

Under the guidance of

Professor Shiv Brat Singh & Dr. Debalay Chakrabarti

DEPARTMENT OF METALLURGICAL AND MATERIALS ENGINEERING INDIAN INSTITUTE OF TECHNOLOGY KHARAGPUR MAY 2016

© 2016 Arka Mandal. All rights reserved.

Dedicated to MY GRANDPARENTS &

TEACHERS

APPROVAL OF THE VIVA-VOCE BOARD

Date: __ /__ /____

Certified

that

the

thesis

entitled

PROCESSING-MICROSTRUCTURE-

MICROTEXTURE-PROPERTY CORRELATION OF DUPLEX STAINLESS STEELS, submitted by ARKA MANDAL to the Indian Institute of Technology, Kharagpur, for the award of the degree Master of Science has been accepted by the external examiners and that the student has successfully defended the thesis in the viva-voce examination held today.

(Member of the DAC)

(Member of the DAC)

Prof. S. B. Singh (Supervisor)

(External Examiner)

(Member of the DAC)

Dr. D. Chakrabarti (Supervisor)

(Chairman)

CERTIFICATE

This is to certify that the thesis entitled Processing-Microstructure-MicrotextureProperty Correlation of Duplex Stainless Steels, submitted by Arka Mandal to Indian Institute of Technology, Kharagpur, is a record of bona fide research work under our supervision and we consider it worthy of consideration for the award of the degree of Master of Science (Research) of the Institute.

__________________________ (Supervisor) Prof. S. B. Singh Professor, Dept. of Metallurgical and Materials Engineering Indian Institute of Technology Kharagpur – 721 302, India

Date:

______________________ (Co-Supervisor) Dr. D. Chakrabarti Associate Professor, Dept. of Metallurgical and Materials Engineering Indian Institute of Technology Kharagpur – 721 302, India

DECLARATION I certify that a. The work contained in the thesis is original and has been done by myself under the general supervision of my supervisors.

b. The work has not been submitted to any other Institute for any degree or diploma. c. I have followed the guidelines provided by the Institute in writing the thesis. d. I have conformed to the norms and guidelines given in the Ethical Code of Conduct of the Institute. e. Whenever I have used materials (data, theoretical analysis, and text) from other sources, I have given due credit to them by citing them in the text of the thesis and giving their details in the references. f. Whenever I have quoted written materials from other sources, I have put them under quotation marks and given due credit to the sources by citing them and giving required details in the references.

(Arka Mandal)

ACKNOWLEDGEMENTS

I would like to thank my supervisors, Prof. S. B. Singh and Dr. D. Chakrabarti for their unstinted support and guidance in this research. Their valuable suggestions, thought provoking ideas and constant encouragement in systematically carrying out experiments, analyses and proper interpretation of results obtained and finally preparation of the thesis are gratefully acknowledged. I am grateful to Head, Dept. of Metallurgy & Materials Engineering (MME), Principal Investigator (P.I.), Steel technology centre (STC) and Chairman, Central Research Facility (CRF) all from IIT Kharagpur for the provision of research facility. I shall remain ever indebted to Shyam Ferro Alloys Ltd. for sponsoring me to have the opportunity to pursue this research work. I would like to thank Mr. Amit Mukherjee (GM, Tech., BRG) for taking interest in my work, providing chemical analysis facility of the steel melts from Durgapur and suggesting studying the industrial aspect of the research work time to time. I take this opportunity to thank Prof. K. K. Ray, Prof. S. Das and all the faculty members of the Department of Metallurgical and Materials Engineering for their valuable suggestions during the course of this investigation. I also wish to thank all the staff members of STC, Melting & casting lab, CRF, CWISS and Dept. of MME, specially to Kamal da, Hemanto da, Prashanta da, Abhipsu da, Basu da, Sudipto, Bishwarup da, Tinku, Nilay for their various types of assistance during experimental work. I would like to acknowledge my departmental seniors, especially Sudipta, Anish, Hasan da, Abhijit da, Snehanshu da, Arya da, Ravi bhaiya for offering me valuable research inputs at different stages, which definitely helped me to get accustomed with the research work. Apart from that, I was lucky to discover many good friends who provided a cozy atmosphere and made my stay at Kharagpur a memorable one. I sincerely thank each one of them, particularly to Ayan, Soumen, Bishu da, Lalu da, Utpal, Abhisek, Arkadeep, Santi da, Sujit da, Monojit da, Arindam da, Kaushik da, Tanmoy, Aritra, Soumi, Sudip etc. It is their friendship and sense of fraternity that sustained me through these years. ‘Masi’ has always been there to serve world class food which is something that should not be left mentioning.

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I cannot find appropriate words to convey my sense of gratefulness to my parents who have done so much unconditionally up to this stage. Thanks are due to my grandparents who were the first teachers of my life. From the very beginning to this stage, all my teachers have been absolutely wizard for me. Lastly, I would prefer to mention my fiancé, Monalisa, for having patience and providing encouragement and support in all my endeavors. Thank you all for everything.

(Arka Mandal)

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LIST OF SYMBOLS

∆V

: Change in volume

A

: Cross-sectional area of the reduced section of tensile specimen, mm2

a, c

: Lattice parameter

A, n

: Constants

AR

: Average aspect ratio (dmax/dmin)

aγ

: Lattice parameter of γ-phase

aδ

: Lattice parameter of δ-phase

b

: Burgers vector

Bs

: Brass texture component

Cr2N

: Needle like Nitride phase, forms in two ranges of temperature (≥1040 °C & 700-950 °C)

Creq

: Chromium equivalent

CrN

: Cubic Nitride phase, forms in the HAZ during welding

d

: Lamellar spacing, mm

D(T)

: Volume diffusion co-efficient

D0

: Constant

dmax

: Maximum grain diameter, µm

dmin

: Minimum grain diameter, µm

dsb

: Sub-grain diameter

Ep

: Upper potential of passive range

Epp

: Lower potential of passive range

G

: Ni, Si, Mo rich intermetallic phase

G

: Shear modulus

GS

: Grain size in ASTM number

H

: Applied field, Oe

Hv

: Hardness in Vickers

HV0.0025

: Vickers hardness measured at 2.5 g load

HV0.01

: Vickers hardness measured at 10 g load

HV1.0

: Vickers hardness measured at 1 kg load

imax

: Maximum current density (in log scale) vii

ipass

: Rate of general corrosion

K

: Partitioning co-efficient

L

: Overall length of tensile specimen, mm

L

: Liquid phase

L1

: Length of the grip section of tensile specimen, mm

m

: Dislocation mobility

M23C6

: Carbide phase, forms rapidly between 650 and 950°C at δ/γ boundaries

M7C3

: Carbide phase, forms within 10 minutes between 950 and 1050°C at δ/γ boundaries

Md30

: The temperature where 50% martensite forms at 30% true strain

Mns

: Type of inclusion, initiation site for pit

ms

: Magnetic saturation, emu/g

n

: Strain hardening exponent

n

: Chromium equivalent / Nickel equivalent

Nieq

: Nickel equivalent

Q

: Activation energy for hot working

Q*

: Activation energy for diffusion

R

: Any random texture

r

: Fillet radius of tensile specimen, mm

R

: Universal gas constant

R

: Mo rich Laves phase (Fe2Mo), precipitates in small quantities between 550 and 650°C

r0, r45 and r90 : Plastic strain ratio in 0°, 45° and 90° to the rolling direction respectively 𝑟̅

: Lankford parameter or normal anisotropy factor

Rm

: Ultimate strength

Rp / Rp0.2

: Proof strength / Proof strength at 0.2% plastic strain

s

: Grain boundary energy

T

: Absolute temperature

t

: Thickness, mm

t, t1, t2

: Time, minute/hour

T, T1, T2

: Temperature, °C viii

T0

: Initial temperature (K)

tf

: Cooling time

Tm

: Absolute melting temperature

v

: Cooling rate (K-s-1)

V

: Volume

W

: Width of the grip section of tensile specimen, mm

w

: Width of the reduced section of tensile specimen, mm

X

: Fractional amount of element in pseudo ternary diagram

α

: BCC ferrite phase

α'

: BCC martensite (stable)

α

: Cr-poor phase, forms due to miscibility gap in Fe-Cr system

α'

: Forms due to miscibility gap in Fe-Cr system between 300 and 525°C

α

: Inverse of the stress associated with power-law breakdown

αBCC, α-fiber

: Characteristic fiber for cold rolled BCC metals ( || RD)

αFCC

: Characteristic fiber for FCC metals ( || ND)

α-SS

: Ferritic stainless steel

β-fiber

: Characteristic fiber for FCC metals

γ

: FCC austenite phase

γ2

: Secondary austenite

γ3

: Tertiary austenite

γSFE

: Stacking fault energy

γ-fiber

: Characteristic fiber for BCC metals ( || ND)

γ-SS

: Austenitic stainless steel

δ

: Content of BCC ferrite phase

δ

: High temperature BCC ferrite phase

ε

: Cu or W rich particles, forms after 100 hours at 500°C

ε

: HCP martensite (metastable)

έ

: Strain rate

εt or ε

: Total strain

θ

: Diffraction angle (half), °

λ

: X-Ray wavelength, Å

ix

π

: Cr and Mo rich phase, forms after isothermal heat treatment at 600°C for several hours

ρ0

: Dislocation density in unrecrystallized material

ρ0c

: Critical dislocation density for dynamic recrystallization

σ

: Cr, Mo rich hard embrittling tetragonal precipitate, forms between 650 and 1000°C

σ

: True stress

τ

: Dislocation line energy

τ

: Heavily faulted needle-like phase, forms between 550 to 650°C on δ/δ boundaries

τ

: Shear stress

φ1, φ and φ2

: Euler angles

χ

: Diffusion distance

χ

: Cr, Mo rich deleterious precipitate, forms between 700 and 900°C

x

LIST OF ABBREVIATIONS AISI

: American Iron and Steel Institute

ASTM

: American Society for Testing and Materials

BCC

: Body-centered cubic

BCT

: Body-centered tetragonal

BF

: Bright field

BSE

: Backscattered electron

CCT

: Continuous cooling transformation

CPT

: Critical pitting temperature

CR

: Cold rolled / Cooling rate

CR_HT850x : Cold rolled sample LD_CR44 heat treated at 850 °C for x hours and then water quenched DRV

: Dynamic recovery

DRX

: Dynamic recrystallization

DSS

: Duplex stainless steel

EBSD

: Electron backscattered diffraction

EBSP

: Electron backscattered patterns

EDS/EDX

: Energy Dispersive Spectroscopy / Energy Dispersive X-ray

FCC

: Face-centered cubic

FN

: Ferrite number

FRT

: Finish rolling temperature

GL

: Gauge length

HAGB

: High angle grain boundary

HAZ

: Heat affected zone

HR

: Hot rolled

HR_HT850x : Hot rolled sample LD_HR50 heat treated at 850 °C for x hours and then water quenched ICx

: Designation for isochronally heat treated SD steel sample (x=serial no.)

IGC

: Intergranular corrosion

IIT Kharagpur : Indian Institute of Technology Kharagpur IMOA

: International Molybdenum Association xi

IPF

: Inverse pole figure

ITx

: Designation for isothermally heat treated SD steel sample (x=serial no.)

K-S

: Kurdjumov-Sachs orientation relationship

LAGB

: Low angle grain boundary

LD

: Lean duplex stainless steel grade melted in air induction furnace

LD_CR44

: AIF steel sample cold rolled by 44% (5 mm in thickness)

LD_HR37

: AIF steel sample hot rolled by 37% (10.7 mm in thickness)

LD_HR50

: AIF steel sample hot rolled by 50% (8.5 mm in thickness)

LD_HT1200 : LD steel sample soaked at 1200 °C for 30 minutes and then water quenched LD_HT8503

: LD steel sample soaked at 850 °C for 3 hours and then water quenched

LD_HT900x

: LD steel sample soaked at 900 °C for x hour(s) and then water quenched

LOM

: Light optical microscopy

MME

: Metallurgical and Materials Engineering

ND

: Normal direction

N-W

: Nishiyama-Wassermann orientation relationships

ODF

: Orientation distribution function

OIM

: Orientation image mapping

PREN

: Pitting resistance equivalent number

P-S

: Petch-Schrader orientation relationship

RD

: Rolling direction

SD

: As-received cold rolled-annealed standard duplex grade steel

SDSS

: Super duplex stainless steel

SE

: Secondary electron

SEM

: Scanning Electron Microscopy

SFE

: Stacking fault energy

SIBM

: Strain induced boundary migration

SRV

: Static recovery

SRX

: Static recrystallization

ST

: Solution treated

TD

: Transverse direction xii

TEL

: Total elongation

TEM

: Transmission Electron Microscopy

TTT

: Time-temperature-transformation

UEL

: Uniform elongation

UNS

: Unified numbering system

UTM

: Universal testing machine

UTS

: Ultimate tensile strength

WRC

: Welding Research Council

XRD

: X-Ray Diffractometer / X-Ray Diffractometry

YS

: Yield strength

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LIST OF TABLES

Table 2.1:

Different grades of DSSs according to their composition and their corresponding PREN value [Alvarez-Armas (2008)].

Table 2.2:

Calculated solidus and liquidus temperatures of duplex stainless steels IC373, IC378 and IC381. Creq and Nieq have been calculated using Hammar and Svensson equation. The measured hardness and volume fraction of austenite of these alloys are also shown [Sharafi (1993)].

Table 3.1:

Chemical composition of the investigated duplex stainless steels (values are in weight %).

Table 3.2:

Processing details of standard duplex (SD) steel samples.

Table 3.3:

Processing details of lean duplex (LD) steel samples.

Table 4.1:

Phase fraction data calculated by Thermo-Calc® for steel LD at 50 ˚C interval (values are in volume %).

Table 4.2:

Phase analysis by light optical microscopic images and X-ray diffraction plots of as-received sample SD and laboratory made as-forged sample LD.

Table 4.3:

Average Vickers micro-hardness and macro-hardness of as-received cold-rolled and annealed steel SD and as-forged steel LD.

Table 4.4:

Chemical analysis of as-received sample SD and as-forged sample LD by Energy Dispersive Spectroscopy (values are in weight %).

Table 4.5:

Volume % of phases in sample SD after cycle – I.

Table 4.6:

Phase fraction (%) of both the phases in SD steel grade (after cycle – II).

Table 5.1:

Average composition of ferrite and austenite phases in SD steel (weight %).

Table 5.2:

Phase fraction by image analysis of Lean DSS (LD) samples isothermally heat treated at 900 ˚C or 850 °C.

Table 5.3:

Aspect ratios of austenite islands in heat treated LD samples LD_HT8503 (850 °C, 3 h) and LD_HT9003 (900 °C, 3 h).

Table 5.4:

Details of image analysis, average aspect ratio of austenite islands and average bulk hardness value of hot rolled samples LD_HR37 and LD_HR50.

Table 5.5:

Microhardness of γ-austenite and δ-ferrite phase in hot rolled samples LD_HR37 and LD_HR50. Indentation load was 25 gm-f and 10 s taken for each reading.

Table 5.6:

Result obtained from image analysis of hot rolled sample LD_HR50 and hot rolled annealed samples HR_HT8500.25 (850 °C, 15 minutes) and HR_HT8503 (850 °C, 3 h).

xiv

Table 5.7:

Results obtained from optical image analysis and hardness measurement of cold rolled sample LD_CR44.

Table 5.8:

Results of image analyses obtained from optical microstructures of cold rolled and annealed samples CR_HT8500.25 (850 °C, 15 minutes) and HR_HT8503 (850 °C, 3 hours).

Table 7.1:

Details of the tensile tested samples: LD_HT8503, LD_HR50, LD_CR44 and CR_HT8503.

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LIST OF FIGURES

Fig. 1.1:

Positioning of Duplex grades - an excellent combination of high strength and corrosion resistance. [Outokumpu (2011)]

Fig. 2.1:

The calculated equilibrium partitioning of (a) carbon and (b) nitrogen between ferrite and austenite in duplex stainless steel alloy IC378 [Sharafi (1993)].

Fig. 2.2:

Effect of alloying elements on the solubility of nitrogen in liquid Fe18%Cr-8%Ni alloys at 1600 °C at 1 atm N2 [Small & Pehlke (1968)].

Fig. 2.3:

Schematic summary of the effects of alloying elements on the anodic polarization curve. It clearly shows that Cr extends the passive range and reduces the rate of general corrosion (ipass) [Sedriks (1985)].

Fig. 2.4:

WRC-1992 diagram [Kotecki & Siewert (1992)]. The FN prediction is only accurate for weld compositions that fall within the bounds of the iso-FN lines (0 to 100 FN) that are drawn on the diagram. The limits of the diagram were determined by the extent of the database and extension of the lines could result in erroneous predictions.

Fig. 2.5:

Calculated equilibrium volume fraction of austenite in duplex stainless steel alloys IC373 and IC381 with Creq/Nieq (n) of 4.09 and 3.04 respectively [Sharafi (1993)].

Fig. 2.6:

(a) Concentration profiles in Fe-Cr-Ni constitution diagram at 70% and 60% Fe. The schematic effect of N addition is shown by darker portion inside the diagram [Schafmeister & Ergang (1939)], (b) Computer calculated ‘isopleth diagram’ with the dotted line indicating the composition of superduplex alloys, e.g. 25Cr-7Ni-4Mo-0.3N [Nilsson (1992)].

Fig. 2.7:

TTT diagrams of duplex stainless steels derived by optical metallography between 600 and 1050 °C and hardness measurements between 300 and 600 °C [Charles (1991)].

Fig. 2.8:

(a) Phase proportions versus heat treatment temperatures and (b) Hardness curve versus heat treatment temperatures. [Martins & Casteletti (2005)]

Fig. 2.9:

(a) Effect of aging treatments on Charpy V impact energy of 2205 duplex stainless steel aged at 400–500 °C (impact energy of as-received material 304 J). (b) Effect of aging treatments on HRc macro-hardness of 2205 duplex stainless steel aged at 400–500 °C (hardness of as-received material HRc 22). [Weng et al. (2004)]

Fig. 2.10:

Continuous cooling diagram from 1080 °C. The hashed area denotes typical HAZ cooling rates, while the two shaded areas depict different sensitivities to intermetallic formation in superduplex alloys [Charles (1991)].

xvi

Fig. 2.11:

Temperature dependence of element partitioning coefficients (K = ferrite/austenite) for a range of duplex steels [Charles (1994)].

Fig. 2.12:

(a) Nitrogen content of ferrite and austenite for different alloys, (b) Temperature dependence of partition coefficients (K) for different grades [Charles (1994), Hertzman et al. (1992)].

Fig. 2.13:

Time-Temperature-Transformation diagram for S32304 steel [Solomon (1982)].

Fig. 2.14:

(a) Coarse sigma precipitates in superduplex plate after 10 minutes at 1000 °C, ×800. Etch: electrolytic sulphuric acid [Chance et al. (1982), Ohmori & Maehara (1984)]. (b) SEM micrograph of γ2 + σ eutectoid. Steel S32750 after 72 hours at 700 °C [Nilsson et al. (1992)].

Fig. 2.15:

Morphology of the sigma phase with respect to the isothermal annealing temperature; (a) 750 °C and (b) 950 °C [Pohl et al. (2007)].

Fig. 2.16:

TEM micrographs showing σ phase and M23C6 carbide particle precipitated at the δ/γ interface, (a) BF image of σ phase and M23C6 carbide; (b) schematic diagram of (a); (c) X-ray diffractogram showing the existence of σ and M23C6 precipitates in 2205 DSS [Chen & Yang (2001)].

Fig. 2.17:

TEM micrographs of Cr2N at: (a) intragranular and (b) intergranular sites [Nilsson (1992)]. The nitrides are found to pin the migrating δ/γ phase boundaries.

Fig. 2.18:

SIBM in the vicinity of the δ/γ interface, complemented by twinning, observed at a strain of 1.3: (a) TEM bright-field image; (b) corresponding schematic interpretation (A and F denote austenite and ferrite respectively, GB indicates the migrating austenite grain boundary) [McKay & Cizek (2001)].

Fig. 2.19:

True stress and true strain curves at different temperatures (a) strain rate = 0.1 s-1, (b) strain rate = 1 s-1. [Mao et al. (2003)]

Fig. 2.20:

Optical micrographs of 1.4362: isothermal deformations at 1180 °C with logarithmic strain of 0.3 at a strain rate of 5 s-1 accumulating to a total strain of (a) εt = 0.3, (b) εt = 1.2 and (c) εt = 2.1. [Herrera et al. (2008)]

Fig. 2.21:

Typical texture components shown in an orientation distribution function (ODF) of (a) BCC and (b) FCC steels [Ul-Haq et al. (1994), Keichel et al. (2003), Pawelski et al. (1978)].

Fig. 2.22:

Hot rolling texture of (a) ferrite and (b) austenite phase in DSS 1.4362 after a total strain of 2.1 at 1180°C and ε = 5s-1. [Herrera et al. (2008)]

Fig. 2.23:

(a) Cold rolled material thickness reduction 80% of a transversal section, SEM–BSE micrograph, (b) X-ray diffractogram of as-received material and cold rolled material (80% thickness reduction). [Baldo & Mészáros (2010)]

Fig. 2.24:

(a) Micrograph of DSS S32304 specimen (t = 1.01 mm, ε = -0.138). Martensite portions inside the austenite island are highlighted by circles. xvii

(b) Magnetization curve of DSS S32304 specimen (t = 0.69 mm, ε = 0.519) [Tavares et al. (2014)]. Fig. 2.25:

SEM micrographs of SAF 2205 aged at 800 °C for 1h as a function of solution treatment temperature; (a) 1050 °C and (b) 1250 °C. [Cho & Lee (2013)]

Fig. 2.26:

φ2 = 45° sections of the ODFs of (a) and (c) ferrite and (b) and (d) austenite of (a) and (b) cold rolled and (c) and (d) cold rolled and annealed at 1200 °C samples. [Hamada et al. (2010)]

Fig. 2.27:

(a) Representative proof stress (Rp0.2) in the rolling and transverse directions, as function of plate thickness and (b) Petch plot relating proof stress (average of 0° and 90° tests) to lamellar spacing d. [Hutchinson et al. (1986)]

Fig. 2.28:

Engineering stress–strain curve of DSS samples after various treatments. [Ghosh et al. (2012)]

Fig. 2.29:

(a) Nominal stress–strain curves at various temperatures and (b) Nominal stress–strain curves with various strain rates at 296 K in the S32101 steels. [Tsuchida et al. (2014)]

Fig. 2.30:

SEM micrographs of fracture surfaces in (a) hot rolled, (b) hot rolledannealed, (c) cold rolled and (d) cold rolled-annealed DSS samples. [Ghosh et al. (2012)]

Fig. 3.1:

Schematic diagram of the thermal schedule used for dilatometry.

Fig. 3.2:

Schematic diagram of isochronal (IC) and isothermal (IT) heat treatment schedules.

Fig. 3.3:

Schematic diagram of DSS strip and samples taken for microstructural, EBSD and tensile property study. t is total thickness.

Fig. 3.4:

Geometry of sub-size tensile specimens.

Fig. 3.5:

Schematic diagram of the region (hashed zone) from where the tensile specimen was chosen for microstructural-microtextural study and fracture analysis. t is total thickness.

Fig. 4.1:

Fe-Cr-Ni pseudo ternary diagrams at (a) 1200 °C, (b) 1100 °C, (c) 1000 °C, (d) 900 °C, (e) 800 °C and (f) 650 °C. (Computed using ThermoCalc®) The black dot indicates the composition of the materials used in this study.

Fig. 4.2:

Phase fraction diagrams of composition (a) SD and (b) LD. (Computed using Thermo-Calc®)

Fig. 4.3:

TTT diagrams of (a) SD and (b) LD steel grades, computed by JMatPro®.

Fig. 4.4:

CCT diagrams of (a) SD and (b) LD steel grades, computed using JMatPro®. Cooling curves for four different cooling rates, i.e., 0.028, 0.28, 2.8 and 28 C-s-1 are presented.

Fig. 4.5:

Positions of investigated steels SD and LD inside the Schneider diagram.

xviii

Fig. 4.6:

Optical microstructures of (a) as-received cold-rolled + annealed SD sample and (b) as-forged LD sample, rolling direction is parallel to the marker.

Fig. 4.7:

Comparative X-ray diffraction plots of as-received cold rolled and annealed SD steel and as-forged LD steel samples.

Fig. 4.8:

Line mapping in a region of sample SD, (a) representative optical image, line mapping of (b) all elements together, (c) Mn (blue), (d) Cr (red), (e) Ni (green) and (f) Mo (violet).

Fig. 4.9:

Dilatation (change in length vs. temperature) curve obtained from cycle – I, material: SD.

Fig. 4.10:

Microstructure of sample SD after Dilatometry Cycle – I.

Fig. 4.11:

(a) Phase fraction diagram of SD, with arrows denoting the change in phase fraction with temperature and (b) the dilatation curve of cycle – I, with arrows denoting the heating and cooling curve.

Fig. 4.12:

(a) The dilatation curve (cycle – II), and (b) microstructure of the SD steel after dilatometry cycle - II; darker austenite (γ) islands, comparatively brighter ferrite (δ) matrix and the brightest sigma (σ) phase are labeled here.

Fig. 4.13:

SEM images of dilatometry cycle – II sample which clearly revealed brighter sigma phase: (a) Secondary Electron (SE) image and (b) Back Scattered Electron (BSE) image of the same region as in (a).

Fig. 4.14:

Deviation of alloy concentration from the mean value in γ, δ and σ phase respectively, material: SD (after dilatometry cycle – II).

Fig. 4.15:

(a) Phase fraction diagram of SD, with arrows denoting the change in phase fraction with temperature and (b) the dilatation curve of cycle – II, with arrows denoting the heating and cooling curve.

Fig. 5.1:

Light optical microstructures of samples heat treated for 30 minutes at (a) 1000 °C (IC1), (b) 1050 °C (IC2), (c) 1100 °C (IC3), (d) 1150 °C (IC4), (e) 1200 °C (IC5), (f) 1250 °C (IC6) and (g) 1200 °C (LD_HT1200) [(a) to (f): SD material and (g): LD material]. The rolling direction is parallel to the marker.

Fig. 5.2:

Amount of δ-ferrite phase after annealing treatment (material: SD and LD) in the temperature range of 1000 – 1250 °C.

Fig. 5.3:

The average contents (weight %) of the important alloying elements in ferrite and austenite phases of the specimens measured using EDX combined with SEM. Four specimens of SD material heat treated for 30 minutes at 1000 °C (IC1), 1050 °C (IC2), 1150 °C (IC4) and 1250 °C (IC6) were arbitrarily chosen for chemical analysis.

Fig. 5.4:

High resolution secondary electron images of SD material sample (IC6) heat treated at 1250 °C for 30 minutes and quenched to room temperature. The Cr2N precipitate along δ/δ grain boundary is pointed in the diagram.

xix

Fig. 5.5:

The EDX datum of a particular point on sample heat treated at 1250 °C for 30 minutes (IC6), (a) spectrum 1 is denoting the point of consideration, (b) the EDX spectrum containing Fe, Cr, Ni, Mo, Mn and Si; the weight % of Cr, Mn, Ni and Mo is given.

Fig. 5.6:

Light optical microstructures of SD samples isothermally heat treated at 900 °C for: (a) 0.25 h (IT1), (b) 0.5 h (IT2), (c) 1 h (IT3), (d) 2 hours (IT4), (e) 5 hours (IT5) and (f) 8 hours (IT6). The samples were quenched in cold water after the treatment. The rolling direction is parallel to the marker.

Fig. 5.7:

Amount of γ, δ and σ phase after ageing treatment at 900 °C for the time range of 0.25 – 8 hours. (Samples obtained from SD material)

Fig. 5.8:

The average contents (weight %) of the important alloying elements in δferrite, γ-austenite and sigma (σ) phases of the SD sample treated at 900 °C for 30 minutes (IT2). The measurement done using EDX combined with SEM. Only one sample (IT2) from SD material was arbitrarily chosen as a representative case.

Fig. 5.9:

Light optical microstructures of isothermally heat treated LD samples: (a) 900 °C, 1 h (LD_HT9001), (b) 900 °C, 3 hours (LD_HT9003), (c) 900 °C, 5 hours (LD_HT9005) and (d) 850 °C, 3 hours (LD_HT8503), finally water quenched. Microstructures were prepared using aqua regia etchant.

Fig. 5.10:

The change in % austenite phase in duplex ferrite + austenite structure of isothermally heat treated samples LD_HT9001 (900 °C, 1 h), LD_HT9003 (900 °C, 3 h) and LD_HT9005 (900 °C, 5 h)) with respect to ageing time at 900 °C.

Fig. 5.11:

Light optical microstructures of hot rolled LD samples (RD-ND plane): (a) LD_HR37 and (b) LD_HR50; hot reduction done by 37% (t = 10.7 mm) and 50% (t = 8.5 mm) with FRT of 862 °C and 520 °C respectively. Rolling direction is parallel to the marker.

Fig. 5.12:

Light optical micrographs of hot rolled and hot rolled + annealed samples: (a) LD_HR50 (b) HR_HT8500.25 (850 °C, 15 minutes) and (c) HR_HT8503 (850 °C, 3 hours). Rolling direction is parallel to the marker.

Fig. 5.13:

Phase maps of hot rolled and hot rolled + annealed samples: (a) LD_HR50 (b) HR_HT8500.25 (850 °C, 15 minutes) and (c) HR_HT8503 (850 °C, 3 hours). Rolling direction is parallel to the marker.

Fig. 5.14:

The comparative grain boundary misorientation graphs of LD_HR50, HR_HT8500.25 (850 °C, 15 minutes) and HR_HT8503 (850 °C, 3 hours) in terms of relative frequency of occurrence – (a) δ-ferrite and (b) γaustenite.

Fig. 5.15:

Graphical representation of local average misorientation angle’s relative frequency of occurrence in sample LD_HR50, HR_HT8500.25 (850 °C, 15 minutes) and HR_HT8503 (850 °C, 3 h) – (a) δ-ferrite and (b) γaustenite. The color code represents the grain in terms of misorientation angle in local misorientation map. xx

Fig. 5.16:

The local misorientation maps of δ-ferrite phases in hot rolled sample (a) LD_HR50 and hot rolled annealed samples: (b) HR_HT8500.25 (850 °C, 15 minutes) and (c) HR_HT8503 (850 °C, 3 h). The color code bar indicates the grain of particular angle of misorientation (in degrees). Rolling direction is parallel to the marker.

Fig. 5.17:

The local misorientation maps of austenite phases in hot rolled sample (a) LD_HR50 and hot rolled annealed samples: (b) HR_HT8500.25 (850 °C, 15 minutes) and (c) HR_HT8503 (850 °C, 3 hours). The color code bar indicates the grain of particular angle of misorientation (in degrees). Rolling direction is parallel to the marker.

Fig. 5.18:

Optical microstructures of cold rolled sample (a) LD_CR44 and cold rolled annealed samples: (b) CR_HT8500.25 and (c) CR_HT8503 (annealed at 850 °C for 15 minutes and 3 hours respectively) and (d) closer view of sample CR_HT8503. Rolling direction is parallel to the marker.

Fig. 5.19:

(a) Secondary electron image of cold rolled sample LD_CR44, lath of αʹmartensite are visible inside γ-islands and (b) αʹ-martensite (volume %) versus thickness reduction (%) curve [Baldo & Mészáros (2010)]; amount of αʹ-martensite is shown after 44% thickness reduction.

Fig. 5.20:

Comparison of the bulk hardness, microhardness of each phase and average aspect ratio of γ-austenite of the annealed samples CR_HT8500.25 (850 °C, 15 minutes) and CR_HT8503 (850 °C, 3 hours) with the cold rolled sample LD_CR44.

Fig. 5.21:

Phase maps of cold rolled sample (a) LD_CR44 and cold rolled + annealed samples: (b) CR_HT8500.25 (850 °C, 15 minutes) and (c) CR_HT8503 (850 °C, 3 h), obtained from EBSD. Rolling direction is parallel to the marker.

Fig. 5.22:

The comparative grain boundary misorientation graphs of LD_CR44, CR_HT8500.25 (850 °C, 15 minutes) and CR_HT8503 (850 °C, 3 hours) in terms of relative frequency of occurrence – (a) ferrite and (b) austenite phase.

Fig. 5.23:

Graphical representation of relative frequency of occurrence of local average misorientation angle in samples LD_CR44, CR_HT8500.25 (850 °C, 15 minutes) and CR_HT8503 (850 °C, 3 hours) – (a) ferrite and (b) austenite phase. The color code represents the grain in terms of misorientation angle in local misorientation map.

Fig. 5.24:

The local misorientation maps of BCC phases in cold rolled and cold rolled + annealed samples: (a) LD_CR44, (b) CR_HT8500.25 (850 °C, 15 minutes) and (c) CR_HT8503 (850 °C, 3 hours). The color code bar indicates the grain of particular angle of misorientation (in degrees).

Fig. 5.25:

The local misorientation maps of austenite phases in cold rolled and cold rolled + annealed samples: (a) LD_CR44, (b) CR_HT8500.25 (850 °C, 15 minutes) and (c) CR_HT8503 (850 °C, 3 hours). The color code bar indicates the grain of particular angle of misorientation (in degrees). xxi

Fig. 6.1:

Inverse pole figure (IPF) color maps of δ-ferrite phase in hot rolled and hot rolled + annealed samples: (a) LD_HR50, (b) HR_HT8500.25 (850 °C, 15 minutes) and (c) HR_HT8503 (850 °C, 3 h).

Fig. 6.2:

Inverse pole figure (IPF) color maps of γ-austenite phase in hot rolled and hot rolled + annealed samples: (a) LD_HR50, (b) HR_HT8500.25 (850 °C, 15 minutes) and (c) HR_HT8503 (850 °C, 3 h).

Fig. 6.3:

φ2=45° sections (Bunge notation) of ODFs of ferrite (BCC) phase in hot rolled and hot rolled + annealed samples: (a) LD_HR50, (b) HR_HT8500.25 (850 °C, 15 minutes) and (c) HR_HT8503 (850 °C, 3 h), (d) standard φ2=45° section for BCC steel showing all texture components [Sk. Md. et al. (2016)].

Fig. 6.4:

ODF intensities of texture components within (a) γ-fiber and (b) α-fiber in BCC phases of hot rolled sample LD_HR50 and hot rolled-annealed samples HR_HT8500.25 and HR_HT8503.

Fig. 6.5:

φ2=0° to 90° sections (Bunge notation) of ODFs of austenite phase in hot rolled and hot rolled + annealed samples: (a) LD_HR50, (b) HR_HT8500.25 (850 °C, 15 minutes) and (c) HR_HT8503 (850 °C, 3 h), (d) standard φ2=0° to 90° sections for FCC steel showing all texture components.

Fig. 6.6:

K-S orientation maps of hot rolled sample LD_HR50 and hot rolled-long term annealed sample HR_HT8503 (850 °C, 3 h), color code represents angle of deviation from K-S.

Fig. 6.7:

Maps of hot rolled sample (a) LD_HR50 and hot rolled-long term annealed sample HR_HT8503, highlighting twin boundaries (red) within austenite grains (yellow).

Fig. 6.8:

Inverse pole figure (IPF) color maps of δ-ferrite phase in cold rolled and cold rolled + annealed samples: (a) LD_CR44, (b) CR_HT8500.25 (850 °C, 15 minutes) and (c) CR_HT8503 (850 °C, 3 h).

Fig. 6.9:

Inverse pole figure color maps of FCC phase in cold rolled sample (a) LD_CR44 and cold rolled-annealed samples (b) CR_HT8500.25 and (c) CR_HT8503, annealed at 850 °C for 15 minutes and 3 hours, respectively.

Fig. 6.10:

φ2=45° sections (Bunge notation) of ODFs of austenite phase in cold rolled and cold rolled + annealed samples: (a) LD_CR44, (b) CR_HT8500.25 (850 °C, 15 minutes) and (c) CR_HT8503 (850 °C, 3 h), (d) IPF map of FCC phase in LD_CR44, highlighting the portion of α′martensite (BCC) in red, (e) standard φ2=45° section for BCC steel showing all texture components [Sk. Md. et al. (2016)].

Fig. 6.11:

φ2 = 0° to 90° sections (Bunge notation) of ODFs of austenite phase in cold rolled and cold rolled + annealed samples: (a) LD_CR44, (b) CR_HT8500.25 (850 °C, 15 minutes) and (c) CR_HT8503 (850 °C, 3 h), (d) standard φ2=0° to 90° sections for FCC steel showing all texture components.

xxii

Fig. 6.12:

K-S orientation maps of cold rolled sample LD_CR44 and cold rolledlong term annealed sample CR_HT8503, color code represents angle of deviation from K-S.

Fig. 6.13:

Maps of cold rolled sample (a) LD_CR44 and cold rolled-long term annealed sample CR_HT8503, highlighting twin boundaries (red) within austenite grains (yellow).

Fig. 7.1:

Engineering stress-strain curves of 1. heat treated sample LD_HT8503, 2. hot rolled sample LD_HR50, 3. cold rolled sample LD_CR44 and 4. cold rolled-annealed sample CR_HT8503.

Fig. 7.2:

SEM micrographs of fracture surfaces of (a) heat treated sample LD_HT8503, (b) hot rolled sample LD_HR50, (c) cold rolled sample LD_CR44 and (d) cold rolled-annealed sample CR_HT8503.

Fig. 7.3:

Light optical images of RD-TD planes of tensile tested LD steel samples: (a) as-forged heat treated LD_HT8503 and (b) cold rolled + annealed CR_HT8503. Tensile loading direction is parallel to the marker.

Fig. 7.4:

SEM images of tensile tested specimens obtained from: (a) as-forged heat treated LD steel sample LD_HT8503 and (b) cold rolled + annealed LD steel sample CR_HT8503. Tensile loading direction is parallel to the marker.

Fig. 7.5:

(a) strain contour map and (b) band contrast map of tensile tested specimen, obtained from cold rolled + annealed LD steel sample CR_HT8503 (done by EBSD). Tensile loading direction is parallel to the marker.

Fig. 7.6:

φ2=45° sections (Bunge notation) of ODFs of ferrite phase in sample LD_HT8503: (a) prior to tensile testing, (b) after tensile testing, and in sample CR_HT8503: (c) prior to tensile testing, (d) after tensile testing.

Fig. 7.7:

φ2 = 0° to 90° sections (Bunge notation) of ODFs of austenite phase in sample LD_HT8503 - (a) prior to tensile testing and (b) after tensile testing.

xxiii

LIST OF FLOW CHARTS

Flow chart 3.1:

Processing route of the as-received material (standard duplex grade, SD). t is thickness.

Flow chart 3.2:

Processing route of laboratory made as-cast steel (lean duplex grade, LD). t is thickness.

xxiv

ABSTRACT

THE effect of processing parameters on the microstructure, microtexture and tensile property of standard and lean duplex stainless steel has been studied extensively. Standard 22Cr5Ni3Mo grade cold rolled sheet and lean duplex grade material made in laboratory air induction furnace were used to perform the experiments. Phase fraction diagrams were calculated using Thermo-Calc® to evaluate the microstructures with respect to temperature. The precipitation kinetics of the intermetallic phases was calculated using JMatPro® software and was compared with the experimental results. Thermal cycles were carried out in dilatometer to elucidate the microstructural evolution in the standard DSS grade. Annealing and ageing treatment has been carried out to understand their effect on phase balance and intermetallic precipitation, especially sigma (σ) and chromium nitride (Cr2N) and, to compare standard and lean DSS material in this respect. Lean DSS material was processed by hot rolling (up to 50%) and cold rolling (~ 44%) followed by annealing at 850 °C for different periods of time to study the effect of these processing parameters on microstructure, microtexture and tensile properties of this material. Cold rolled samples revealed strain induced martensite transformation within γ-islands. Electron back scattered diffraction (EBSD) technique has been employed to study the evolution of microstructure during annealing; changes were studied in terms of grain boundary misorientation and local average misorientation. Inverse pole figure (IPF) maps and orientation distribution functions (ODF) were also plotted to investigate the microtextural evolution during annealing. Cold rolled-annealed material revealed suitable microstructure and texture for formability. Tensile properties of the samples having different microstructures were evaluated and correlated with the processing parameters. Cold rolledannealed material showed optimum strength-ductility property. Fracture surfaces of the tensile tested specimens were analyzed under scanning electron microscope (SEM). Samples were also collected from broken tensile specimens in order to study the microstructural and microtextural changes during static tensile loading.

xxv

xxvi

CONTENTS Page No. Title Page Certificate of Approval Certificate Declaration Acknowledgements List of Symbols List of Abbreviations List of Tables List of Figures List of Flow Charts Abstract Contents Chapter 1

i ii iii iv v vii xi xiv xvi xxiv xxv xxvii

Introduction

1-4

1.1

General background

2

1.2

Objectives of the present work

4

Chapter 2

Literature Review

5-46

2.1

Introduction

6

2.2

Types of stainless steels

6

2.2.1

Transformable / Martensitic stainless steels

6

2.2.2

Ferritic stainless steels

7

2.2.3

Austenitic stainless steels

7

2.2.4

Duplex stainless steels (DSSs)

8

2.3

Physical metallurgy of DSSs

9

2.3.1

Effect of alloying elements

9

a. Interstitial elements

9

b. Substitutional elements

11

c. Chromium equivalent (Creq) and Nickel equivalent (Nieq)

13

d. Phase balance prediction

14

e. Effect of Creq/Nieq ratio on equilibrium volume fraction of γ

15

f. Fe-Cr-Ni system

16

xxvii

Page No. 2.3.2

2.3.3

2.4

2.5 Chapter 3

Heat treatment

17

a. Temperatures above 1050 °C

18

b. The 600-1050 °C nose

18

c. The 300-600 °C nose

19

Element partitioning and intermetallic phases

22

a. Element partitioning

22

b. Characteristic and morphology of precipitates

23

Deformation behavior of DSSs

29

2.4.1

Hot deformation characteristics

29

2.4.2

Cold deformation characteristics

37

Mechanical property

41

Experimental Details

47-57

3.1

Material

48

3.2

Microstructural evolution

48

3.3

Processing

49

3.3.1

Heat treatment

50

3.3.2

Hot rolling

51

3.3.3

Cold rolling

52

3.3.4

Stress relieving annealing

52

3.4

Characterization

53

3.4.1

53

3.4.2

3.4.3

Microstructural characterization a. Sample preparation

54

b. Image analysis

54

c. Hardness measurement

54

d. Lattice parameter determination

54

e. Chemical analysis

54

Microtextural characterization

55

a. Grain boundary misorientation

55

b. Crystallographic orientation

56

Mechanical testing

xxviii

56

Page No. Chapter 4

Phase Diagram and Microstructural Evolution

59-80

4.1

Introduction

60

4.2

Calculation on Phase transformation

60

4.2.1

Fe-Cr-Ni pseudo-ternary diagrams

60

4.2.2

Phase fraction diagrams

62

4.2.3

Time-Temperature-Transformation (TTT) curves

64

4.2.4

Continuous Cooling Transformation (CCT) curves

65

4.2.5

Schneider-diagram

67

4.3

Microstructural characterization

69

4.4

Microstructural evolution

73

4.4.1

Dilatometric study

73

4.4.2

Analysis of dilatation

79

4.5 Chapter 5

Summary and Conclusion Effect of Processing Parameters on Microstructure

80 81-117

5.1

Introduction

82

5.2

Heat treatment

82

5.2.1

Annealing treatment

82

5.2.2

Ageing treatment

88

5.3

5.4

5.5 Chapter 6

Hot rolled and hot rolled-annealed steels

95

5.3.1

Hot deformation

95

5.3.2

Effect of annealing on hot rolled microstructure

97

Cold rolled and cold rolled-annealed steels

104

5.4.1

Cold deformation

104

5.4.2

Effect of annealing on cold rolled microstructure

106

Summary and Conclusion Effect of Processing Parameters on Microtexture

115 119-139

6.1

Introduction

120

6.2

Hot rolled and hot rolled-annealed steels

120

6.3

Cold rolled and cold rolled-annealed steels

130

6.4

Summary and Conclusion

138

xxix

Page No. Chapter 7

Effect of Processing Parameters on Tensile Property

141-154

7.1

Introduction

142

7.2

Tensile property of DSSs

142

7.3

Fracture analysis

144

7.4

Microstructural and microtextural changes during tensile loading

146

7.4.1

Microstructural observation

146

7.4.2

Microtextural observation

149

7.5 Chapter 8

Summary and Conclusion Conclusions and Scope for Further Work

153 155-161

8.1

Conclusions

156

8.2

Scope for further work

161

List of References

163

xxx

Chapter 1

INTRODUCTION

1.1 General background 1.2 Objectives of the present work

Chapter 1

Introduction

1.1 General background Duplex stainless steels (DSSs) were introduced in the market of stainless steel in the '40s as they showed less propensity towards intergranular corrosion (IGC) in various corrosive media. This, combined with the superior strength that could be as high as 3 times the yield strength of austenitic stainless steels [Payson (1932), Hochmann et al. (1977)], made DSSs became popular, especially in pipe manufacturing for chemical processing, oil and gas exploration, refinery, marine environments, pulp and paper industries etc. The positioning of duplex grades with respect to austenitic grade materials in terms of strength and corrosion resistance is shown in Figure 1.1.

Fig. 1.1: Positioning of Duplex grades - an excellent combination of high strength and corrosion resistance. [Outokumpu (2011)] In addition, DSSs possess good resistance to fatigue and corrosion fatigue, good abrasion and erosion resistance, high energy absorption, low thermal expansion, good weldability and formability [Outokumpu (2013)]. These attributes of DSSs attracted many researchers to explore the structure-property correlation under different processing parameters. The effect of alloying elements on phase balance and intermetallic precipitation was extensively studied by many scientists. Recent development in duplex chemistry has been made by adding 0.2-0.3 wt% N, as it can replace costly Ni and act as austenite stabilizer. Cr and Mo are found to extend passive range [Hashimoto et al. (1979), Sedriks (1985)], but, higher use can lead to deleterious intermetallic precipitation, such as sigma (σ) or chi (χ) [Roscoe & Gradwell (1986)]. Previously proposed phase prediction-Schaeffler diagram was modified several times 2

Chapter 1

Introduction

and finally WRC-1992 diagram was developed by Kotecki & Siewert (1992), which was in good agreement with the practical data available for weld metal [Nassau (1982)]. Later, isopleth diagram was designed by Nilsson & Liu (1988) for a given composition. Several authors worked on annealing and ageing treatment with varying temperatures and reported various intermetallic precipitations in different temperature regions. Amongst these phases, sigma is the most detrimental as corrosion originates from these sites. Many papers were published on sigma phase formation and its different morphologies [Redjaimia et al. (1991), Strutt et al. (1986), Maehara et al. (1983), Pohl et al. (2007)]. Rolling of duplex stainless steels has always been challenging as the phases (γ and δ) act differently under rolling load and cracks can initiate from the interphase boundaries [Tsuge (1991)]. The recrystallization mechanisms are also different for the two phases. Ferrite recrystallizes through extended dynamic recovery or continuous dynamic recrystallization [McKay & Cizek (2001), Gourdet & Montheillet (2000)], whereas, austenite undergoes conventional dynamic recrystallization, i.e., nucleation and growth [Botella et al. (1996), Jiménez et al. (1998)]. Recent findings are related to development of microtexture in DSSs. Hot rolling develops typical rolling texture of BCC metals in ferrite [Raabe & Lücke (1992,1993)] and recrystallization texture in austenite as it undergoes DRX which inhibits the advancement of typical rolling texture [Herrera et al. (2008), Humphreys et al. (2001)]. Austenite phase undergoes strain induced martensitic transformation under cold deformation [Milad et al. (2008), Baeva et al. (1995), Baldo & Mészáros (2010)] and instability of austenite increases with increased amount of N or Mn [Zhang et al. (2009)]. Magnetization saturation increased with increasing cold deformation due to 𝛾 → 𝛼′-martensite transformation [Tavares et al. (2014)]. Plastic deformation can influence intermetallic precipitation too [Júnior et al. (2011), Choi et al. (2011)]. Recent investigation reveals that cold deformation results in typical cold rolling texture of BCC and FCC metals, i.e.¸ strong α-fiber in δ-ferrite and MS-Brass with strongest component {110} in γ-austenite [Fargas et al. (2008), Keichel et al. (2003)]. Cold rolled-annealed sample has been reported to possess maximum UTS×TEL value and its fracture surface is essentially of ductile nature [Ghosh et al. (2012)]. Nevertheless, overall effect of processing parameters (i.e., heat treatment, hot rolling and cold rolling) on microstructure and microtexture of both standard and lean

3

Chapter 1

Introduction

duplex grade steels have not been thoroughly researched and compared. Accordingly, the present work was undertaken to elucidate the effect of processing parameters and, microstructural and microtextural changes associated with them in standard and lean DSS grade materials.

1.2 Objectives of the present work The major objectives of the investigation as emerged from the literature survey are listed below: 1. To examine the effect of processing parameters, i.e., heat treatment, hot rolling, cold rolling on microstructure and microtexture of duplex grade steels. 2. To validate the calculated thermodynamic and kinetic data with the experimental results for the standard and lean duplex grade steels. 3. To elucidate the change in tensile property and fracture surface associated with the change in processing parameter. 4. To investigate the microstructural and microtextural changes under tensile loading and to study the instability of austenite against plastic deformation.

4

Chapter 2

LITERATURE REVIEW

2.1 Introduction 2.2 Types of stainless steels 2.3 Physical metallurgy of DSSs 2.4 Deformation behavior of DSSs 2.5 Mechanical property

Chapter 2

Literature Review

2.1 Introduction Stainless steels were introduced during the first decades of the twentieth century in the UK and Germany. Martensitic and ferritic Fe-Cr steels were the first in this category. But, due to the ease of production and fabrication (specially welding), austenitic Fe-CrNi steel grades were well received and became popular at a rapid pace. However, C level could not be refined below 0.08 weight% due to the constraints in refining technologies. That made the steel sensitive to grain boundary carbide precipitation during heat treatment and welding. These grain boundaries acted as the most potential sites for intergranular corrosion. A duplex alloy is defined as a two-phase structure where both the phases are present in significant quantities. The steel was first introduced in cast form [Bain & Giffith (1927)]. After several castings with different compositions, in 1933, an error during the melting of an 18%Cr-9%Ni-2.5%Mo grade at the Firminy works of the J. Holtzer Company, France, led to the production of a 20%Cr-8%Ni-2.5%Mo steel. Subsequent analysis [Hochmann (1950)] of the casting found it had a high volume fraction of ferrite in an austenitic matrix and was not sensitive to intergranular corrosion (IGC) in various corrosive media. This was a significant discovery, as the high carbon austenitic grades of the time tended to form a continuous chromium carbide network, leading to rapid corrosion in the surrounding chromium depleted zones., This observation, combined with the superior strength over the austenitic grades [Payson (1932), Hochmann et al. (1977)] encouraged French patents to be issued in 1935 and 1937 [French patent (1935, 1937)].

2.2 Types of stainless steels According to the microstructure, the stainless steels are conventionally categorized as follows: 2.2.1 Transformable / Martensitic stainless steels Martensitic stainless steels usually contain 12-17 wt.% Cr and 0.1-1.5 wt.% C. The lower carbon varieties still have high hardenability, which renders them air hardenable so that tempering is necessary to produce useful combinations of strength, toughness and ductility. The common type of steel, 11-13 wt.% Cr [Briggs & Parker (1965)] is alloyed with Mo, W, V and Nb to increase its tempering resistance, which often necessitates the addition of up to 3 wt.% Ni to eliminate ferrite from the structure, which would otherwise decrease the potential strength. AISI steel type 403, 410, 416, 420, 6

Chapter 2

Literature Review

440 are some well established martensitic steels market grades. These steels are mainly used as cutlery and tool materials. 2.2.2 Ferritic stainless steels These steels usually contain 15-30 wt.% Cr with alloying additions of Mo, Ti or Nb and are fully ferritic up to the melting point. They have relatively low work hardening rates compared with the austenitic stainless steels, but higher yield stresses due to their bodycentred cubic (BCC) crystal structure. They exhibit a ductile-brittle fracture transition which often leads to a lack of toughness, but they have the advantage over the austenitic steels of showing virtual immunity to chloride induced stress corrosion cracking. Whilst having good corrosion resistance, which can be increased by higher Cr contents and Mo additions, they are generally less corrosion resistant than the austenitic steels and also somewhat less formable. They substitute however for the more expensive austenitic materials in the less severe corrosion environments and where formability requirements are less stringent [Brandis et al. (1974)]. Typical applications of fully ferritic steels are in domestic, catering, architectural and decorative uses. They are used widely in the chemical, food, transportation and automobile industries. AISI type 405, 430, 444, 446 are some typical ferritic market grade stainless steels. 2.2.3 Austenitic stainless steels These types of stainless steels are the most widely used, comprising 70-80 wt. % of stainless production [Leach (1982)]. They usually contain 16-25 wt.% Cr, 7-20 wt.% Ni and current developments have decreased the carbon content to well below 0.03 wt. %. They are frequently alloyed with Mo, Ti and Nb to promote creep resistance and, with other than lowest carbon contents, stabilization against intergranular corrosion. The austenitic steels have lower proof stresses but higher work hardening rates than the ferritic stainless steels, and are more easily welded and cold formed. Their toughness is high, exhibiting no ductile-brittle fracture transition and their exceptional corrosion resistance is mainly marred by their susceptibility to chloride induced stress corrosion cracking. Whilst being essentially austenitic at all temperatures, the austenitic steels may contain (delta) ferrite, a small amount of which is detrimental to hot workability [Bywater & Gladman (1976)]. Anyway some delta ferrite is essential for good weldability. The applications of austenitic stainless steels are widespread in the chemical processing, food, constructional, power plant and other industries in which

7

Chapter 2

Literature Review

their corrosion resistance, cryogenic properties, creep and oxidation resistance are of predominant importance. AISI type 304, 304L, 305, 308, 309, 316, 316L steels are established market grades of austenitic stainless steel. 2.2.4 Duplex stainless steels (DSSs) Duplex stainless steels are a relatively new class in the stainless steels family, although they have been known for almost 45 years [Truman (1979)]. They comprise more or less equal amounts of austenite and ferrite. The appropriate structure can be produced by a wide range of compositions, usually relatively lean in Ni and varying in Cr from 18-30 wt.%, with additions of Mo and Mn, and in some cases stabilization by Nb or Ti. The addition of N is a recent development. Whilst there is partitioning of alloying elements between the ferrite (Cr and Mo) and austenite (Ni and Mn), it is stated that this is not excessive [Redmond & Miska (1982)] and has no serious effect on corrosion resistance. The properties of duplex steels are somewhere between the properties of austenitic and ferritic steels. Duplex steels have high resistance to the stress corrosion cracking and to chloride ions attack. These steels are weldable and formable and possess high strength. The application of modern duplex stainless steels is in plate form and as tubes in the paper industry, as heat exchangers and for various uses in chemical engineering. A major use, however, is as a casting alloy or forgings in pumps, valves and general marine engineering. UNS S32205, S31803, S32304 are the standard 22% Cr grades and UNS S32520, S32750 are some of super duplex grades. In recent years, some leaner alloy grades, e.g. S32101 have been developed to cut off the cost of expensive alloying elements (Ni, Mo) with an insubstantial loss of mechanical and corrosion resistance properties. The success of the 22050 grade led to the development of an entire family of duplex alloys, which range in corrosion resistance depending on their alloy content [IMOA (2014)]. It is common to define the corrosion resistance of duplex grades by their pitting resistance equivalence number [Gunn (1997)] (PREN) as defined by Eq. 2.1: PREN = %Cr + 3.3 × %Mo + 16 × %N

(2.1)

PREN is derived from an empirical relationship and can take several forms. This number does not provide an absolute value for corrosion resistance and is not applicable in all environments; it gives an overview of the expected resistance to pitting corrosion 8

Chapter 2

Literature Review

in an aqueous chloride solution. Nevertheless, the most widely employed for duplex alloys is given above. According to the PREN, the DSSs have been categorized in three groups and are shown along with their chemical constitution in Table 2.1. Table 2.1: Different grades of DSSs according to their composition and their corresponding PREN value [Alvarez-Armas (2008)]. Grade

UNS

C

Cr

Ni

Mo

W

Cu

N

PREN

Lean Duplex

S32101

0.03

21.5

1.5

0.3

-

-

0.22

25

S32304

0.02

23.0

4.0

0.3

-

0.3

0.10

25

Standard Duplex

S31803

0.02

22.0

5.5

3.0

-

-

0.17

35

S32205

0.03

22.5

5.8

3.2

-

-

0.17

36

Superduplex

S32750

0.02

25.0

7.0

4.0

-

0.5

0.27

43

S32760

0.03

25.0

7.0

3.5

0.6

0.5

0.25

42

2.3 Physical metallurgy of DSSs The overall physical metallurgy of Duplex Stainless Steels (DSSs) is mostly concerned with the following effects and mechanisms associated with them. 2.3.1 Effect of alloying elements It is generally accepted that the favorable properties of DSSs can be achieved for phase balances in the range of 30 – 70% ferrite and rest austenite. However, DSSs are most commonly considered to have roughly equal amounts of ferrite and austenite. The interactions of the major alloying elements, particularly chromium, molybdenum, nitrogen and nickel are quite complex. To achieve a stable duplex structure care must be taken to obtain the correct level of each of the alloying elements. In duplex stainless steel, a decreased amount of Ni is added (as compared to austenitic stainless steel - 8%) to have a ferritic-austenitic duplex structure. a. Interstitial elements Carbon and Nitrogen are used as major interstitial alloying elements in DSSs. Their concentration plays significant role in duplex alloys. Carbon Duplex stainless steels usually contain less than 0.08 wt.% carbon (mostly in range 0.025-0.035 wt.%). It is an austenite (γ) stabilizer and increases the maximum solubility of Cr in austenite, which is beneficial for corrosion resistance property [Gordon (1977)]. Anyway, to ensure good hot workability carbon content should be restricted to a maximum of 0.030 wt%. The low carbon concentration inevitably suppresses carbide 9

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precipitation and further reduces the susceptibility of duplex stainless steel to localized corrosion in the as-welded condition [Roscoe & Gradwell (1986)]. Figure 2.1a shows the calculated equilibrium partitioning of carbon between ferrite and austenite in alloy IC378 [Sharafi (1993)]. At 1350 °C the carbon content of austenite is about five times as much as of ferrite. This illustrates that although the density of austenite is higher than ferrite the solubility of interstitial carbon atoms is greater in austenite. This is due to the empty spaces in the FCC structure are fewer but larger for interstitial carbon atoms. As the temperature decreases, the carbon content of ferrite approaches that for austenite and the remaining ferrite with high carbon content becomes unstable and ready to transform to austenite.

Fig. 2.1: The calculated equilibrium partitioning of (a) carbon and (b) nitrogen between ferrite and austenite in duplex stainless steel alloy IC378 [Sharafi (1993)]. Nitrogen Duplex stainless steels can have up to 0.3 wt.% nitrogen which is an austenite stabiliser. Nitrogen is more soluble than carbon in the Fe-Cr-Ni system. The addition of 3 to 5 wt.% Mn increases nitrogen solubility in duplex stainless steel castings to 0.35 wt.% with acceptable internal soundness [Sharafi (1993)]. Previous report has shown that high levels of nitrogen will necessitate high levels of manganese [Chance et al. (1982)]. It is often used as a cheap nickel substitute in austenitic and duplex stainless steels [Leslie (1981)]. As far as the corrosion properties are concerned, one of the most significant improvements in the localized corrosion performance of duplex stainless steels is obtained by increasing nitrogen levels in the region of 0.2-0.3 wt.% [Chance et al. (1982)]. Furthermore, the nitrogen in solid solution increases the strength of the 10

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alloy [Roscoe & Gradwell (1986)]. Figure 2.1b shows the calculated equilibrium concentrations of nitrogen in ferrite and austenite in alloy IC378 [Sharafi (1993)]. Nitrogen has also been noted to increase the crevice corrosion resistance. Researchers have proposed [Clayton & Martin (1988)] that this is due to nitrogen altering the crevice solution chemistry or by segregating to the surface, which is in keeping with the mechanism for enhanced pitting resistance [Newman & Shahrabi (1987)]. The effect of other alloying elements on the solubility of nitrogen has also been discussed previously and it suggests that more the concentration of alloying elements more the nitrogen solubility (Figure 2.2).

Fig. 2.2: Effect of alloying elements on the solubility of nitrogen in liquid Fe-18%Cr8%Ni alloys at 1600 °C at 1 atm N2 [Small & Pehlke (1968)]. b. Substitutional elements Amongst all substitutional alloying elements, Cr, Mo, Ni and Mn are the prime in all aspects. This is not only because they contribute in phase balance, but also they play vital roll individually to enhance mechanical or corrosion resistance property. Chromium The main advantage of adding chromium to steel is to improve the localized corrosion resistance by the formation of a passive chromium-rich oxy-hydroxide film [Hashimoto et al. (1979)]. Electrochemically this is achieved by extending the passive range [Sedriks (1985)] as shown in Figure 2.3. Nevertheless, there is an upper limit in Cr 11

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content to be added, as higher content leads to deleterious intermetallic precipitation, such as sigma (σ). These precipitates result in substantial loss in ductility, toughness and corrosion resistance property.

Fig. 2.3: Schematic summary of the effects of alloying elements on the anodic polarization curve. It clearly shows that Cr extends the passive range and reduces the rate of general corrosion (ipass) [Sedriks (1985)]. Molybdenum Molybdenum is well named for its beneficial impact on pitting and crevice corrosion in chloride medium. It plays the same role as Cr by extending the passive range and reducing the corrosion current density (imax) in the active range (as depicted in Figure 2.3). The mechanism by which molybdenum increases the pitting resistance of an alloy has been examined by a number of workers [Sedriks (1979), Newman & Franz (1984), Olefjord et al. (1983)], and has been found to suppress active sites via formation of an oxy-hydroxide or molybdate ion [Halada et al. (1995)]. However, an upper limit for Mo addition is quoted of about 4% [Roscoe & Gradwell (1986)] as Mo also enhances intermetallic precipitation, such as chi (χ), sigma (σ). Nickel To balance the volume % of both phases, the level of nickel addition to a given duplex alloy will depend primarily on the chromium content. Excessive nickel content results 12

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into high austenite content and more elements partitioning in remaining ferrite. As an effect, chances of intermetallic precipitation from ferrite increase in supportive temperature range [Goldschmidt (1951)]. In brief, nickel does have direct effect on corrosion properties, for instance moving Ep in the noble direction and reducing ipass (Fig. 2.3), and yet it appears that the main role of nickel is to control phase balance and element partitioning. Manganese Although Mn has BCC lattice, it has been proven as an austenitic stabilizer in most of the cases [Schaeffler (1949), DeLong et al. (1956)], however, exception of this phenomenon has also been noticed by the previous workers [Chance et al. (1982), Guiraldenq (1967)]. It does not have pronounced effect on phase balance and thus was not included in Ni equivalent by many researchers. Mn is added to increase abrasion and wear resistance [Soulignac & Dupoiron (1990)] and tensile properties without any loss in ductility [Roscoe et al. (1984)]. Nonetheless, high Mn content may increase the temperature range of detrimental phase formation [Roscoe et al. (1984)]. Mn is also known to have negative impact on critical pitting temperature (CPT) by forming MnS inclusions which can act as initiation sites for pits [Sedriks (1983)]. In recent years, Ni is substituted by Mn as an austenite stabilizer in lean duplex grade, mainly owing to cost cutting. c. Chromium equivalent (Creq) and Nickel equivalent (Nieq) Chromium and other elements stabilize ferrite, although the effect of different elements varies. Equation has been derived by previous researcher [Kotecki & Siewert (1992)] to quantify elemental effects (the so-called chromium equivalents, Creq) and the equation is as follows: Creq = %Cr + %Mo + 0.7 × %Nb + 1.5 × %Si

(2.2)

The ferrite formers contribute in chromium equivalent and the increase in this value results in higher volume % of ferrite in final ‘duplex’ microstructure. Amongst these, Cr and Mo are paramount as they play the most significant role in properties of DSSs. Counter to the ferrite stabilizing effect of Cr, Mo, Si and Nb, there is another group of elements which stabilize austenite: Nieq = %Ni + 35 × %C + 20 × %N + 0.5 × %Mn + 0.25 × %Cu 13

(2.3)

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The alloying elements contributing in nickel equivalent are necessarily austenite stabilizers. In order to maintain about 40% to 60% ferrite and balance austenite, the ferrite stabilizing elements, Eq. 2.2, must balance with the austenite stabilizers, Eq. 2.3. Ni and Mn are the most influential in phase balance amongst all the austenite formers. d. Phase balance prediction The Fe-Cr-Ni system has been examined in several ways since first published data by Bain and Griffith. The Schaeffler diagram [Schaeffler (1949)] is an empirical determination of the microstructures of weld metal that result from welding different compositions. This diagram has been used for many years to predict the cast or weld metal microstructures in conventional austenitic and stainless steels and to optimize base metal and filler compositions [Gunn (1997)]. It became apparent that the Schaeffler diagram did not anticipate duplex microstructures very well [Hoffmeister & Mundt (1978)] and so other diagrams and relationships were recommended [Noble & Gooch (1986), Skuin & Kreyssing (1978)], culminating in the WRC-1992 diagram (Figure 2.4) [Kotecki & Siewert (1992)].

Fig. 2.4: WRC-1992 diagram [Kotecki & Siewert (1992)]. The FN prediction is only accurate for weld compositions that fall within the bounds of the iso-FN lines (0 to 100 FN) that are drawn on the diagram. The limits of the diagram were determined by the extent of the database and extension of the lines could result in erroneous predictions. The WRC diagram proffers reasonable agreement in many practical exercises [Nassau (1982)], in spite of the fact that for relatively low and high nitrogen contents (0.26%N), FN values are over- and under-estimated respectively 14

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[Zhang et al. (1996)]. This may indicate that a variable factor for nitrogen is required, as proposed by DeLong in 1956 [DeLong et al. (1956)] or that cooling rate should be considered [Yasuda et al. (1994)]. Nevertheless, the WRC-1992 diagram is still recommended for first approximation of weld metal content. e. Effect of Creq/Nieq ratio on equilibrium volume fraction of γ Thermodynamic calculations have shown that the equilibrium volume fraction of austenite decreases as the Creq/Nieq ratio of steel increases. Table 2.2 shows the Creq/Nieq of hot rolled duplex stainless steels IC373, IC378 and IC381 calculated using existing equation [Hammar & Svensson (1979)]. Figure 2.5 shows that the equilibrium volume fraction of austenite of IC373 with Creq/Nieq=4.09 is much lower than of IC381 with Creq/Nieq=3.04 [Sharafi (1993)]. The calculations were done for the temperature range where only austenite and ferrite are thermodynamically stable. Table 2.2: Calculated solidus and liquidus temperatures of duplex stainless steels IC373, IC378 and IC381. Creq and Nieq have been calculated using Hammar and Svensson equation. The measured hardness and volume fraction of austenite of these alloys are also shown [Sharafi (1993)]. Steel Grade IC373

4.09

Liquidus Temp. °C 1484

313

Austenite Volume fraction 0.3

IC378

3.12

1474

1384

281

0.57

IC381

3.04

1470

1382

273

0.63

Creq/Nieq

Solidus Temp. °C 1382

Hardness

Fig. 2.5: Calculated equilibrium volume fraction of austenite in duplex stainless steel alloys IC373 and IC381 with Creq/Nieq (n) of 4.09 and 3.04 respectively [Sharafi (1993)]. 15

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f. Fe-Cr-Ni system Firstly, schematic form of Fe-Cr-Ni diagram was represented considering 70% Fe level [Schafmeister & Ergang (1939)] and is depicted in Figure 2.6a. The phase proportions and their respective chemical compositions are indicated for an alloy analysis and annealing temperature, with more influence of nitrogen in phase balance than Cr or Mo at high temperature. As shown in Figure 2.2 [Small & Pehlke (1968)], the addition of 0.25%N to a 25 % Cr alloy produces a ferrite volume fraction of about 50% at 1250°C, compared to nearly 80% ferrite with 0.18 % N. However, it is inconvenient to prognosticate the microstructure of a duplex alloy from simplified diagrams, due to the effects of other alloying elements, which significantly modify the phase fields.

Fig. 2.6: (a) Concentration profiles in Fe-Cr-Ni constitution diagram at 70% and 60% Fe. The schematic effect of N addition is shown by darker portion inside the diagram [Schafmeister & Ergang (1939)], (b) Computer calculated ‘isopleth diagram’ with the dotted line indicating the composition of superduplex alloys, e.g. 25Cr-7Ni-4Mo-0.3N [Nilsson (1992)]. 16

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One approach [Sundman et al. (1985)] has been to construct isothermal ternary sections of Fe-Cr-Ni-Mo-N using a computer program called ThermocalCTM which computes phase equilibria, over a range of temperatures. This can then be employed [Nilsson (1992)] to present an isopleth diagram for a given composition (Fig. 2.6b). This approach has been found to be in good agreement with experimental data. 2.3.2 Heat treatment Plenty of structural changes develop in DSSs during isothermal or continuous cooling. Most of the transformations are associated with δ-ferrite as diffusion is at least 100 times faster in ferrite than in austenite due to its less compact lattice structure. Most deleterious phases contain Cr and Mo which were in ferrite phase prior to intermetallic precipitation. Ferrite phase becomes supersaturated with these alloying elements as temperature gets lowered, leading to intermetallic precipitation. Time-Temperature-Transformation (TTT) diagrams Time-Temperature-Transformation (TTT) diagrams, generated by isothermal heat treatments followed by quenching, are seldom used to elucidate the susceptibility of different DSS grades to embrittlement and are shown in Figure 2.7 [Charles (1991)]. In this diagram, within the 600-1050 °C temperature range, the curves are based on optical microscope observations, while in the range 300-600 °C, they are determined from hardness measurements.

Fig. 2.7: TTT diagrams of duplex stainless steels derived by optical metallography between 600 and 1050 °C and hardness measurements between 300 and 600 °C [Charles (1991)]. 17

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The diagrams are conservative and do not acknowledge the change in mechanical or corrosion resistance properties as soon as the boundaries indicated here, are reached. Nevertheless, an effect on properties is witnessed practically before the boundaries are reached. a. Temperatures above 1050 °C Unlike steel grades with full austenitization, DSSs solidify completely in ferritic field as γ-austenite formation gets restricted (closed γ-field) owing to high Cr or (Cr+Mo) concentration in DSS chemistry [Bhadeshia & Honeycombe (2011)]. Solid state transformation 𝛿 → 𝛾 takes place upon cooling (Fig. 2.6a). This reaction is a reversible process and 𝛾 → 𝛿 transformation takes place upon heating above 1050 °C. Reduction in partition coefficients of substitutional elements occurs with increasing temperature, i.e., K tends to unity (Fig. 2.7). Additionally, ferrite gets enriched in interstitial elements like C and N. High nitrogen content (0.25-0.40 wt.%) may lead to increase in stability of duplex structure. At this concentration level, the phase ratio is more or less 1, whereas, grades with 0.2 wt.% N roughly have 80-85% ferrite content [Gunn (1997)]. Like other steel grades, grains in DSSs can be made equiaxed by prolonged heat treatment at 1100-1200 °C. Acicular grains with Widmanstätten type structure may form when subjected to intermediate cooling rate from annealing temperature. A mixture of coarse and fine austenite grains can be obtained after step quenching [Charles (1991)]. b. The 600-1050 °C nose Intermetallic precipitation strongly depends on sample chemistry, specifically, concentration of sigma forming elements such as Cr, Mo at this temperature range. Thus, S32304 grade, leaner with Cr and Mo, found to be the least prone to intermetallic precipitation (Fig. 2.14). This grade requires at least 10 to 20 hours to begin the precipitation below 900 °C and thus can be allowed to solution anneal below 1000 °C. On the other hand, grade S32205/S31803 (Mo content 2.5-3.5 wt.%) experiences more propensities towards intermetallic precipitation than the leaner grade S32304 (Mo content is very low). Due to this, higher solution annealing temperature (>1000 °C) should be employed for these steels. Newly developed super duplex grades (SDSS) display the greatest tendency for intermetallic precipitation as these are mostly enriched with Cr, Mo and W. Nonetheless, the severity of precipitation kinetics of these SDSSs 18

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can be equivalent to super-austenitic or super-ferritic stainless steels, at worst [Mancia et al. (1987), Cao et al. (1987), Charles et al. (1987), Perteneder et al. (1986)]. Problem arises when higher gauge materials (>60mm thickness) [Cozar et al. (1994)] are subjected to cooling. Open bore holes accelerate the cooling, particularly in midsection. Redissolution of the precipitates while annealing requires higher temperature for SDSSs than for comparatively leaner standard duplex grades. Similar high temperatures are required for weld joints while solution annealing treatment [Liljas (1994)]. This temperature range is the most dangerous as the most deleterious phases precipitate at this range. In recent years, investigation sheds light on the temperature dependence on phase balance and change in hardness value associated with it [Martins & Casteletti (2005)]. Figure 2.8 depicts the results obtained by Martins & Casteletti (2005). The graphs clearly show that increase in annealing temperature results in less amount of sigma phase and, therefore, a sharp decrease in Brinell hardness value above 950 °C has been witnessed in standard duplex grade. c. The 300-600 °C nose At this temperature range, the leaner alloy grade S32304 is less sensitive to hardening than grades with high Mo content. A much shorter time is needed for grades S32205 or S31803 to start the hardening, whereas, the incubation time is much lengthier in case of S32304 [Charles (1991)]. SDSSs experience the most proneness towards hardening in wide temperature range and possess shortest incubation period due to higher Cr and Mo contents and, if present, copper additions. In this temperature range, spinodal decomposition of ferrite takes place and Cr-poor α and Cr-rich αʹ is formed. This embrittling phenomenon is commonly known as ‘475 °C embrittlement’. Ni, Si and Mo rich G-phase also develops in this domain but in very small amount [Iacoviello et al. (2005)]. These intermetallic precipitations are detrimental in terms of impact toughness. Workers showed that the impact energy dropped drastically within 100 minutes of ageing at 475 °C, whereas, it remained almost unchanged even after 1000 minutes of ageing at 400 °C in S32205 DSS grade [Weng et al. (2004)]. The corresponding hardness value increased with ageing time and the rate of increase is the maximum in sample aged at 475 °C. The findings made by Weng et al. (2004) are shown in Figure 2.9a (impact energy) and 2.9b (hardness).

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Fig. 2.8: (a) Phase proportions versus heat treatment temperatures and (b) Hardness curve versus heat treatment temperatures. [Martins & Casteletti (2005)]

Fig. 2.9: (a) Effect of aging treatments on Charpy V impact energy of 2205 duplex stainless steel aged at 400–500 °C (impact energy of as-received material 304 J). (b) Effect of aging treatments on HRc macro-hardness of 2205 duplex stainless steel aged at 400–500 °C (hardness of as-received material HRc 22). [Weng et al. (2004)] Since, diffusion rate is quite slow at this temperature range, the formation of αʹ or G-phase is possible only if the steel is subjected to long-term thermal ageing. Formation of these precipitates is practically impossible in case of lean DSSs.

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Continuous Cooling Transformation (CCT) diagrams Continuous cooling takes place in industrial practices. At temperature near the solvus, the nucleation rate of precipitate is slow and growth rate is fast, whereas, near the nose of the transformation curve, the case is reverse. Thus, it is hard to circumvent sigma precipitation while cooling thick sections and redissolution of these intermetallics is needed by further annealing treatment at sufficiently higher temperature. However, the slow nucleation rate at higher temperature and slow growth rate at lower temperature make it comparatively easy to avoid sigma precipitation [Gunn (1997)]. Researchers performed the continuous cooling operation at different cooling rates in grade S32520 [Charles (1991)]. The continuous cooling diagram is shown in Figure 2.10. A significant shift in the position of the precipitation ‘nose’ between the TTT (Fig. 2.7) and the CCT (Fig. 2.10) can be witnessed [Charles (1991), Josefsson et al. (1991)]. It is possible to have precipitation free ‘duplex’ structure even if cooling rate as slow as 2000 °C-h-1 is employed from solution annealing temperature of about 1080 °C.

Fig. 2.10: Continuous cooling diagram from 1080 °C. The hashed area denotes typical HAZ cooling rates, while the two shaded areas depict different sensitivities to intermetallic formation in superduplex alloys [Charles (1991)]. Nevertheless, fast quenching rate should be opted in case of higher annealing temperature, since, Cr2N nitride can precipitate at this temperature range (as mentioned in section 2.4.3). The cooling rates measured in the heat affected zones (HAZs) of welds are generally faster (Fig. 2.10) than can be achieved during quenching, which negates intermetallic formation [Gunn (1997)].

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2.3.3 Element partitioning and intermetallic phases Depending upon the solution treatment temperature and cooling rate from that temperature, different secondary phases precipitate at various temperature regions owing to element partitioning. a. Element partitioning The phase diagram shown in Fig. 2.6a illustrates that the ferrite and austenite will have different compositions depending on the temperature, further illustrated by experimental data [Charles 1994] in Fig. 2.7. Several workers [Strutt et al. (1986), Herbsleb & Schwaab (1982), Solomon (1982), Miyuki et al. (1983), Lardon et al. (1987)] have studied the composition of the phases and resultant partition coefficients in solution annealed product and have demonstrated that, with the exception of nitrogen, the partition coefficients for a given element do not vary significantly between alloys, in spite of the wide range of compositions investigated. This is due to the solubility limits for these elements not being exceeded, for the concentrations and annealing temperatures concerned. Nevertheless, for the most highly alloyed grades (superduplex steels), it has been proposed [Charles (1994)] that the partition coefficients tend more towards unity, apparently due to the use of higher solution annealing temperatures (Fig. 2.11).

Fig. 2.11: Temperature dependence of element partitioning coefficients (K = ferrite/austenite) for a range of duplex steels [Charles (1994)].

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The partition of nitrogen in alloys is observed to vary substantially with chemical composition. However, the nitrogen partitioning coefficient is also controlled by temperature. During a solution annealing treatment, even though the solubility of nitrogen in ferrite may enhance a little, the volume fraction of austenite decreases markedly. This leads to enrichment of nitrogen in the remaining austenite [Charles (1994)], which is depicted in Figure 2.12a. This results in increase in nitrogen partitioning, i.e., K tends to shift from unity [Charles (1994), Hertzman et al. (1992)]. Figure 2.12b depicts this result. The results obtained from these diagrams are in well agreement with the computed ThermocalCTM results [Hertzman et al. (1992)].

Fig. 2.12: (a) Nitrogen content of ferrite and austenite for different alloys, (b) Temperature dependence of partition coefficients (K) for different grades [Charles (1994), Hertzman et al. (1992)]. b. Characteristic and morphology of precipitates The Duplex stainless steels are prone to many intermetallic phases’ precipitation in different stages of heat treatment. The Time-Temperature-Transformation diagram of Duplex grade S32304 has been shown in Figure 2.13. Among these phases sigma (σ) is the most deleterious one. Apart from that, chi (χ) phase is also likely to precipitate. But, due to the thermodynamical instability, this intermetallic remains at a small fraction if not been seen at all in final microstructures of DSSs. Carbides (M23C6 and M7C3) and nitrides (Cr2N and CrN) are also witnessed in case of high Cr and high (Cr + N) bearing DSSs respectively when exposed to high temperature annealing [Gunn (1997)]. 23

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Fig. 2.13: Time-Temperature-Transformation diagram for S32304 steel [Solomon (1982)]. Study on nitride, sigma phase and concomitant γ-phase formation was undergone in this present work and is depicted in this thesis. Hence, literature survey of only these phases is described in the next sub-sections. Sigma (σ) [Fe-Cr-Mo – tetragonal] The deleterious sigma phase is a hard, embrittling, Cr and Mo rich precipitate. Sigma is known to precipitate at temperature region between 650-1000 °C with nose at around 900 °C as seen in Fig. 2.13. This precipitate affects corrosion resistance and impact toughness properties extensively. Some researchers suggested that sigma phase precipitates on the pre-existing M23C6 particle [Goldsmith (1967), Blenkinsop & Nutting (1967), Weiss & Stickler (1972)]. Some evidences of this mechanism of development of sigma phase were witnessed in DSSs above 750 °C temperature [Redjaimia et al. (1991), Wang et al. (1991)]. Sigma nuclei can grow into course plates or lamellar eutectoid σ + γ2 [Strutt et al. (1986), Chance et al. (1982), Ohmori & Maehara (1984)] aggregates, as depicted in Figure 2.14a and 2.14b respectively. This secondary austenite (γ2) was sometimes named as tertiary austenite (γ3) by other workers [Pohl et al. (2007)]. The elements Cr, Mo, Si, Mn and Ni are known to have positive effect on sigma phase formation [Chance et al. (1982), Roscoe et al. (1984)]. However, Ni reduces the equilibrium volume fraction of sigma in final microstructure [Maehara et al. (1983)]. Ni stabilizes austenite and as an effect of this Cr concentration 24

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increases in ferrite phase which, in turn, increases the propensity towards sigma phase formation. Finally, 10% plastic strain results in decrease in required time for sigma phase formation by an order of magnitude [Maehara et al. (1983)].

Fig. 2.14: (a) Coarse sigma precipitates in superduplex plate after 10 minutes at 1000 °C, ×800. Etch: electrolytic sulphuric acid [Chance et al. (1982), Ohmori & Maehara (1984)]. (b) SEM micrograph of γ2 + σ eutectoid. Steel S32750 after 72 hours at 700 °C [Nilsson et al. (1992)]. As the time temperature transformation diagram shows (Fig. 2.13), the fastest precipitation rate for sigma phase can be found between 850 and 900 °C [Roscoe et al. (1984), Strutt et al. (1986)]. The morphology of the sigma phase precipitation changes according to the precipitation temperature. At lower precipitation temperatures of 750 °C, a coral-like or lacy [Martins & Casteletti (2009)] structure of sigma phase can be observed (Fig. 2.15a).

Fig. 2.15: Morphology of the sigma phase with respect to the isothermal annealing temperature; (a) 750 °C and (b) 950 °C [Pohl et al. (2007)].

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The amount of sigma nucleus at the initiating stage of the precipitation is rather high, depending on the shorter diffusion distances at lower precipitation temperatures as reported by Pohl et al. (2007). Therefore, lower diffusion velocity leads to higher local supersaturation and results in a higher density of precipitations. A different precipitation mechanism can be witnessed at higher temperatures of 950 °C (Fig. 2.15b). The sigma phase is larger and denser at these temperatures and the connection between single sigma crystals is minimal, resulting from a lower nucleation formation force but a high diffusion rate at elevated temperatures [Pohl et al. (2007)]. Further research on the effects of solution treatment and continuous cooling on σ-phase precipitation was carried out in 2205 DSS grade [Chen & Yang (2001)]. TEM microscopy revealed well retained M23C6 within σ phase (Figure 2.16a). The conservation of M23C6 particle within σ phase had been reported previously in the austenitic stainless steel by Lai et al. (1981), who attributed this phenomenon to the low carbon solubility of in σ phase.

Fig. 2.16: TEM micrographs showing σ phase and M23C6 carbide particle precipitated at the δ/γ interface, (a) BF image of σ phase and M23C6 carbide; (b) schematic diagram of (a); (c) X-ray diffractogram showing the existence of σ and M23C6 precipitates in 2205 DSS [Chen & Yang (2001)]. The precipitation behavior of intermetallic during continuous cooling is suggested that primarily, small M23C6 particles were precipitated at the δ/γ interface and within the δ ferrite grain interior by conventional nucleation mechanism; then σ 26

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phase nucleated preferentially at the high energy δ/γ interface as well as on the preprecipitated M23C6 particles and grew into the δ ferrite. Meanwhile, the metastable δ ferrite adjacent to the δ/γ interface transformed to new δ phase. As an effect of this, the M23C6 particles pre-formed at δ/γ interface were surrounded within the grown γ region, as can be seen schematically in Fig. 2.16b. In addition to the major diffractions from δ ferrite and γ austenite, minor diffraction peaks from σ phase and M23C6 carbide could be identified in X-Ray diffractometry after solution annealed the sample at 1080 °C and cooled at 0.1 °C-s-1. The X-ray diffractogram was reported by Chen & Yang (2001) and is shown in Figure 2.16c. Maehara et al. (1983) substantially studied δ-γ DSSs with composition ranging: 19.4-25.5 wt% Cr, 4.9-9.7 wt% Ni and 0.03-3.0 wt% Mo. Their studies concluded Cr, Mo enrichments in σ-phase and Ni enrichment in γ-austenite. Depending upon the concentration of sigma forming elements such as Cr and Mo in ferrite and δ/γ ratio before precipitation, σ-phase mainly transforms through eutectoid decomposition of 𝛿 → 𝜎 + 𝛾3 . The increase in Cr and Mo concentration results in increase in total volume % of σ precipitates. The increase in amount of Ni content as well as the decrease in solution annealing temperature brings about higher rate of acceleration in σ precipitation and vice versa. However, Ni reduces the total amount of σ-precipitation and solution annealing temperature does not have significant effect on this. Nitrides Cr2N and CrN [cubic] Nitrogen is added to DSSs to stabilize austenite. Nitrogen is more soluble in austenite than in ferrite and gets partitioned in former phase [Kolts et al. (1983)]. When DSSs are subjected to high temperature annealing at above 1040 °C, ferrite phase becomes thermodynamically more stable than austenite and 𝛾 → 𝛿 transformation takes place (Figure 2.6a). At this temperature the solubility of nitrogen is quite high in ferrite, but solubility drastically decreases as cooling takes place, i.e., low temperature ferrite, super saturated with nitrogen appears. This leads to intragranular precipitation of needle-like Cr2N and is shown in Figure 2.17a. DSSs experience more propensities towards nitride precipitation if higher solution annealing temperature [Roscoe et al. (1984)] and/or rapid cooling rate from that temperature is employed [Nilsson & Liu (1988)]. Isothermal exposure in temperature range 700-950 °C promotes intergranular

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Cr2N along δ/δ grain boundaries as thin plates on sub-grain boundaries, triple points, inclusions and on dislocation arrays (Figure 2.17b).

Fig. 2.17: TEM micrographs of Cr2N at: (a) intragranular and (b) intergranular sites [Nilsson (1992)]. The nitrides are found to pin the migrating δ/γ phase boundaries. Cubic CrN develops in DSSs after welding operation and it is seen more in heat affected zone (HAZ) [Kajimura et al. (1991), Hertzman et al. (1986)]. CrN precipitates adjacent to the Cr2N particles and appears in film-like or tiny platelet-like shape. The orientation relation between CrN and δ-ferrite matrix was found to be [110]CrN || [111]δ and (001)CrN || (11̅0)δ. Both Cr2N and CrN nitrides are enriched in Cr, Fe, V and Mo [Liao (2001)]. Chen & Yang (2002) researched thoroughly on Cr2N precipitation and after analyzing the electron diffraction patterns and isometric projections, they suggested that Cr2N rods were predominantly precipitated in Petch-Schrader (P-S) orientation relationship with δ matrix [Chen & Yang (2002)]. In order to understand the diffusibilities of chromium and nitrogen atoms in δ ferrite, the diffusion distances (χ) of chromium and nitrogen atoms in δ phase during a continuous cooling were calculated by researchers and represented by Equation 2.4 [Kobayashi et al. (2000)]. 𝑡

𝜒 = √∫0 𝑓 𝐷(𝑇)𝑑𝑡

(2.4)

in which, tf is the cooling time. The volume diffusion coefficient D(T)can be expressed by Equation 2.5.

𝐷(𝑇) = 𝐷0 𝑒𝑥𝑝(−

28

𝑄∗ 𝑅𝑇

)

(2.5)

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The values of D0 and Q*have been discovered by other group of researchers [Maruzen (1993)] and those are 1.13×10-6 m2-s-1 and 83.0(1 - 14.03/T) kJ-mol-1 for nitrogen, and 2.3×10-4 m2-s-1 and 239 kJ- mol-1 for chromium, respectively. The temperature T can be represented by Equation 2.6.

𝑇 = 𝑇0 − 𝑣𝑡

(2.6)

Liao et al. calculated the diffusion distance of Cr and N atoms in standard DSS grade while cooling from 1000 °C to 600 °C. They suggested that the nitrogen atoms travel a distance many times larger than the Cr atoms even at very slow cooling rate. As a result, the recovery of chromium depletion adjacent to the nitride precipitates is burdensome, whereas, the diffusion of nitrogen atoms in order to precipitate CrN nitrite is quite favorable. Another group of researchers worked on microstructural characterization and the effect of phase transformations on toughness of the UNS S31803 duplex stainless steel aged at 850 °C. It concluded that ageing at 850 °C for >6 minutes resulted in chromium nitride precipitation, which was found to cause loss of toughness property in DSSs [Zucato et al. (2002)]. Other than these two deleterious intermetallics, there are some other forms of detrimental phases appear during different stages of heat treatment depending upon the composition of DSSs. According to severity, those phases are Chi (χ) phase (Fe36Cr12Mo10 – BCC-αMn) [Rideout et al. (1951)], Carbides M23C6 and M7C3 [Solomon (1982)], Alpha prime (α') [Williams (1958)], G-phase [Auger et al. (1990)], R-phase [Rideout et al. (1951)], π-phase [Nilsson & Liu (1991)], τ-phase [Redjaimia et al. (1991)], Cu-rich epsilon (ε) phase [Charles (1991)] etc.

2.4 Deformation behavior of DSSs Mechanical metallurgy involves mechanical working of metals and alloys in solid state. Testing mechanical properties, establishing relations between these properties with processing parameter and predicting materials performance during service are also the important areas that come under the purview of mechanical metallurgy. 2.4.1 Hot deformation characteristics The DSSs contain about 50% ferrite phase in their microstructures and, thus, can be rolled/formed at ease at above 950 °C. Super-plastic behavior i.e. an elongation even of several 100% before failure has been witnessed when deformed above 900 °C [Osada

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et al. (1987)]. A very low strain rate, temperature of at least 0.6Tm (Tm is the absolute melting temperature) and very fine structure are essential to obtain superplasticity [Gunn (1997)]. Some recent developments are made on the hot workability of DSSs. A review paper was published by Barteri et al. (2000) on hot forming of duplex stainless steels. A comparative study among DSS, austenitic (γ) SS and ferritic (α) SS was carried out in terms of yield strength and corrosion resistance properties. The DSSs possess high strength which is the result of mutual strengthening effect of α and γ in the microstructures. But, the decrease in strength is greater than both the single phase stainless steels with increasing temperature as ferritic transformation (𝛾 → 𝛼) takes place at elevated temperature. Evidences suggested strain partitioning preferentially to the softer ferrite phase when present in low volume fraction under low strains. Exception of this conclusion was drawn when greater capability in withstanding substructural changes in α than in γ-phases has been found. The α-phase grows well defined subgrains when a low strain is applied, but at high έ, some large misorientations presumably result from deformation band rotation. The γ-phase develops substructures of much higher dislocation density in smaller more highly misoriented cells than in αphase and undergoes discontinuous dynamic recrystallization (DRX) or conventional DRX at critical strains much higher than in simple γ-SS due to the absence of nucleation sites at internal grain boundaries within the γ particles. The DRX in the γ-phase is commonly the cause of the peak in the flow curve but does not instigate DRX in αphase. As annealing progresses, the α-phase undergoes substantial static recovery before recrystallizing, whereas, the γ-phase particles recrystallized readily to single crystals traversed by annealing twins [Barteri et al. (2000)]. DSSs can be processed in many ways, but, problems like edge cracks or inappropriate surface finish can arise while deformation due to different response of the two phases under hot working condition [Tsuge (1991)]. In order to improve the poor workability [Nilsson et al. (1992), Ahlblom & Sandström (1982)], study was carried out to derive constitutive relation for the flow stress under hot working conditions of two different DSSs. At high temperatures, the flow behavior of polycrystalline materials can be described by the classical hyperbolic sine relation [Sellars & Tegarts (1966), Garofalo (1963), Jonas et al. (1969)] and is represented in Equation 2.7. −𝑄

𝜀̇ = 𝐴(𝑆𝑖𝑛ℎ 𝛼𝜎)𝑛 𝑒𝑥𝑝 ( 30

𝑅𝑇

)

(2.7)

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where, A and n are the constants α is the inverse of the stress associated with power-law breakdown Q is the activation energy for hot working 𝜀̇ and σ are the true strain rate and true stress involved, respectively T is the temperature and R is the universal gas constant The parameters A, α, n and Q obtained in this way can be compared with some data reported in the literature [Paúl et al. (1993), Al-Jouni et al. (1983), Maehara (1985)]. However, the experimentally obtained values somewhat differ from those reported in the literature. At high temperatures and different strains, the grain morphology remains almost same as the original microstructure, exclusive of the δ/γ interfaces which were serrated. At low temperatures, regardless of strain rate, the morphology changes drastically as the interfaces become very sharp and ferrite grains seem to deform under heavy shear stresses [Cabrera et al. (2003)]. Austenite can undergo dynamic recrystallization and promote super-plasticity [Botella et al. (1996), Jiménez et al. (1998)]. Although, some authors reported that dynamic recrystallization of austenite was restricted in these alloy grades [Al-Jouni et al. (1983), Iza-Mendia et al. (1998), Iza-Mendia (1999)]. Report suggested that DRX nucleation was commonly termed as strain induced boundary migration (SIBM) in single phase austenitic stainless steel [Humphreys (1991)]. In case of DSSs, nucleation mechanism was frequently detected in the δ/γ interface, as shown in Figure 2.18. It was witnessed that SIBM was commonly accompanied by the multiplication of twins at the migrating boundary. Multiplication of twins occurred either by decreasing the energy of migrating boundary [Sample et al. (1987)] or by increasing its mobility [McQueen (1993)]. Except for SIBM using preexisting austenite high angle boundaries, some DRX nuclei were also formed by recovery process of sub-grains [McKay & Cizek (2001)]. Ferrite softening was a result of intense DRV, complemented by a gradual increase in mean misorientation angle across sub-grain boundaries during straining. This process often entitled as ‘extended DRV’ instead of ‘continuous DRX’ in the literature [Gourdet & Montheillet (2000)]. The processes of migration and merging of low-angle dislocation walls tend to form a complex equilibrium network of small and large angle boundaries [Belyakov et al. (1993)].

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Fig. 2.18: SIBM in the vicinity of the δ/γ interface, complemented by twinning, observed at a strain of 1.3: (a) TEM bright-field image; (b) corresponding schematic interpretation (A and F denote austenite and ferrite respectively, GB indicates the migrating austenite grain boundary) [McKay & Cizek (2001)]. Some other group of researchers elaborately worked on as-cast high Mn steel to study the stress-strain behavior under hot deformation. True stress-true strain curves obtained by those experiments are shown in Figure 2.19. The initial rapid rise of stress is allied with the increase in dislocation density and formation of low angle sub-grain boundaries owing to continuous work hardening and dynamic recovery (DRV). Accumulation and annihilation of dislocations get to a dynamic balance as the stress reaches to a saturation value σss. Eventually, the local dislocation density gets sufficiently high that it can yield driving force for nuclei formation of dynamic recrystallization (DRX) [Mao et al. (2003)].

Fig. 2.19: True stress and true strain curves at different temperatures (a) strain rate = 0.1 s-1, (b) strain rate = 1 s-1. [Mao et al. (2003)]

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The peak strain can be described by Equation 2.8 given by Roberts (1984):

𝜌0 ≥ 𝑐𝑜𝑛𝑠𝑡 (

𝑠𝜀̇ 𝐛𝑑𝑠𝑏 𝑚𝜏2

) = 𝜌0𝑐

(2.8)

where, 𝜌0 is the dislocation density in unrecrystallized material 𝜌0𝑐 is the critical dislocation density for DRX 𝑠 is the grain boundary energy b is the Burgers vector 𝜏 is the dislocation line energy dsb is the sub-grain diameter m is the dislocation mobility and 𝜀̇ is the strain rate This equation confirmed the role of grain boundaries on dynamic recrystallization. The activation energy for deformation was found 480 kJ-mol-1, which is higher than those obtained for α-SS or γ-SS [Cizek & Wynne (1997)]. Increasing temperature and decreasing strain rate promote dynamic recrystallization [Mao et al. (2003)]. In order to study the effect of δ-ferrite co-existence on hot deformation and recrystallization of austenite, workers performed hot torsion tests on AISI 304 γ-SS and 2205 DSS. It is suggested that the DRX fraction as well as DRX grain size is quite larger in case of AISI 304 steel than for 2205 steel under same hot working conditions. The major restoration mechanism in austenitic stainless steels is conventional DRX, i.e., bulging of initial grain boundaries, whereas, restoration takes place through continuous DRX, i.e., coalescence of sub-boundaries in duplex stainless steels [Dehghan-Manshadi & Hodgson (2008)]. Some authors worked on the methods for improving hot plasticity in terms of microstructure and chemical composition. It was discussed that the hot plasticity of DSSs can be enhanced by optimizing the sulphur content and stabilizing the holding temperature or decreasing the temperature variation range before hot deformation [Gang et al. (2008)]. Important conclusions came out when extensive study on microstructure and texture of hot rolled DSS was published together. The optical images of hot rolled structures obtained from the reported result are shown in Figure 2.20 [Herrera et al. (2008)].

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Fig. 2.20: Optical micrographs of 1.4362: isothermal deformations at 1180 °C with logarithmic strain of 0.3 at a strain rate of 5 s-1 accumulating to a total strain of (a) εt = 0.3, (b) εt = 1.2 and (c) εt = 2.1. [Herrera et al. (2008)] The volume fraction of phases remained almost constant for all values of strain; however, 𝛾 → 𝛿 transformation took place while hot rolling. At εt = 0.3, the sub-grains of ferrite evinced a recovered structure, whereas, austenite exhibited inhomogeneous and deformed structure with no evidence of DRV or recrystallization. Annealing twins formed within γ grains during soaking treatment. At εt = 1.2, the ferrite became more elongated and larger sub-grains formed through static recovery (SRV). Fragmented austenite grains were also present, but, instead of recovery, internal boundaries formed by DRX and annealing twins were attributed by static recrystallization (SRX) during the inter-pass time. At εt = 2.1, ferrite exhibited fully recovered structure, whereas, austenite grains were more of internal boundaries and annealing twins due to DRX and SRX [Herrera et al. (2008)]. Further study suggested that low temperature hot deformation could result into serrated δ/γ interface and these could act like the initiation site for microcracks within ferrite [Hölscher et al. (1991)]. It is stated that ferrite recrystallized through DRV as the stacking fault energy is quite higher [Humphreys & Hatherly (2004)]. On the other hand, austenite possesses low stacking fault energy (γSFE ~ 21 mJ-m-2) [Schramm & Reed (1975)]. Therefore, DRV becomes sluggish in austenite and increase in dislocation density facilitates the DRX [Humphreys & Hatherly (2004)]. Duprez et al. (2002) suggested that the crystallographic texture evolution of both phases in all stages of hot rolling is similar to that observed in individual FCC and BCC metals [Raabe (2003), Ul-Haq et al. (1994), Keichel et al. (2003), Pawelski et al. (1978)]. Figure 2.21 depicts typical texture components of BCC and FCC steels in an orientation distribution function. It is noted that the austenite texture mainly consists of β-fiber (characteristic fiber for FCC metals) and a weak cube component, whereas, ferrite exhibits a similar 34

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cold rolling texture with αBCC and γ-fibers [Duprez et al. (2002), Keichel et al. (2003), Pawelski et al. (1978)]. Herrera et al. (2008) studied on the δ/γ interface inside the ferrite matrix with Kurdjumov-Sachs (K-S) or Nishiyama-Wassermann (N-W) orientation relationships. Ferrite revealed strong solidification texture, mainly || ND, whereas, austenite possessed of weak texture component prior to deformation. In case of ferrite, the solidification texture disappears with the applied strain and a weak (111)[112] component on the γ-fiber is formed. The hot deformation results into typical BCC cold rolling texture [Raabe & Lücke (1992,1993)] in DSSs. The weak cube component {100} under small strain transforms to rotated cube component {100} under large strain. This component has a propensity towards recovery instead of recrystallization [Ul-Haq et al. (1994)] due to its small in-grain orientation gradients. The weak (111)[112] component was known not only as a deformation component, but also a recrystallization component of the {112} αBCC-fiber on the ferrite [Duprez et al. (2002)].

Fig. 2.21: Typical texture components shown in an orientation distribution function (ODF) of (a) BCC and (b) FCC steels [Ul-Haq et al. (1994), Keichel et al. (2003), Pawelski et al. (1978)]. On the other hand, ferrite consisted of rotated cube {100}, Goss {110} and (110)[11̅1] components. The αFCC fiber, || ND, was developed with a combination of Goss {011} and Brass {011} components as the 35

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strain increased to 1.2 and rotated cube component disappeared in course of the deformation process. At larger strain (εt = 2.1), the Brass component became the strongest with a combination of weak cube component {100} and || RD β-fibers. The αFCC-fiber and β-fiber with main components at Brass {011}, Copper {112} and S {123}, could be spotted as FCC plane strain rolling texture [Raabe et al. (2002), Humphreys et al. (2001)]. The cube component is sometimes denoted as recrystallization texture in FCC metals [Pawelski et al. (1978), Raabe et al. (2002), Humphreys et al. (2001)]. It is believed that during hot rolling, FCC phase undergoes DRX or SRX which inhibits the advancement of typical rolling texture. Figure 2.22 is showing the hot rolling texture (ODF) of both phases under large strain (εt = 2.1) [Herrera et al. (2008)]. A recently published journal suggested that during prolonged annealing of warm-rolled DSS, the strength of RD-fiber, especially of {112} component increased in ferrite with the expense of {001} component. In case of austenite, annealing revealed strong BS texture component and the retention of rolling texture. Nevertheless, prolonged annealing treatment weakened the BS texture and the maximum intensity came in between G and BS along the α-fiber. Specifically, this component was {110} and φ1, φ and φ2 values were 20°, 45° and 0°, respectively [Zaid Ahmed & Bhattacharjee (2015)].

Fig. 2.22: Hot rolling texture of (a) ferrite and (b) austenite phase in DSS 1.4362 after a total strain of 2.1 at 1180°C and ε = 5s-1. [Herrera et al. (2008)] 36

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2.4.2 Cold deformation characteristics Rolling has always been a major area of attraction for the researchers as the two phases behave differently under mechanical or thermo-mechanical treatment. Since, the rolling load is even higher in case of cold rolling; chance of crack generation from the interface is even more. Many workers have researched on hot deformation and aspects related to it, but, study on cold deformation of DSSs is not extensive. Study on cold rolled (20% thickness reduction) and cold rolled-annealed samples were carried out in late ’90s. Results suggested that austenite phase exhibited higher driving force for recrystallization than ferrite owing to more strain hardening of the former phase. Hardness and dislocation density values of austenite phase were quite larger (434 HV0.0025 and 5.77×1011, respectively) compared to the ferrite phase (358 HV0.0025 and 2.32×1011, respectively) [Reick et al. (1998)].

An old published work suggested

about high strength steel with ferrite and martensite as its microstructural constituents [Hayden & Floreen (1970)]. A recent study highlights ferritic-martensitic steel by cold rolling the ferrite-austenite DSSs. The process is governed by strain induced martensitic transformation of austenite within ferrite matrix. Many authors worked on this mechanism of transformation in AISI 301, 304 and 316 austenitic stainless steels [Tavares et al. (2008), Mangonon & Thomas (1970), Mészáros & Prohászka (2005), Choi & Jin (1997), Seetharaman & Krishnan (1981), Güler et al. (2007)] and standard DSSs [Tavares et al. (2006)]. In recent years, lean duplex grade has been developed with low Ni content and increased Mn or sometimes N content [Zhang et al. (2009)]. This results in higher instability of austenite phase, i.e., strain induced martensite formation becomes easier under cold rolling or heat treatment. Two types of martensite can form in this process: paramagnetic ε-martensite (hcp) and ferromagnetic αʹmartensite (bcc). Since, ε-martensite is metastable compared to αʹ-martensite, ε phase forms before αʹ phase and disappears with formation of latter phase under heavy strain. Based on the microstructural observation, the sequence of transformation is understood in order 𝛾 → 𝜀 + 𝛼ʹ [Milad et al. (2008), Baeva et al. (1995)]. A strong grain refinement has been witnessed in cold rolled steels and shown in Figure 2.23a [Baldo & Mészáros (2010)]. Qualitative or quantitative analysis of strain induced martensite is not possible through light optical imaging and this left scope for XRD analysis for identification of αʹ-martensite. Figure 2.23b depicts the difference between XRD plots of sample without cold deformation and 80% cold deformed sample [Baldo & Mészáros (2010)]. The cold rolled sample did not show αʹ and δ separately as these two phases present 37

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same reflections due to their same lattice structure, i.e., BCC unit cell. The peak of γ completely disappears after heavy cold deformation with the increase in 𝛿 + 𝛼ʹ peaks. This clearly concludes that the austenite undergoes deformation induced transformation to αʹ-martensite in course of the process.

Fig. 2.23: (a) Cold rolled material thickness reduction 80% of a transversal section, SEM–BSE micrograph, (b) X-ray diffractogram of as-received material and cold rolled material (80% thickness reduction). [Baldo & Mészáros (2010)]

Fig. 2.24: (a) Micrograph of DSS S32304 specimen (t = 1.01 mm, ε = -0.138). Martensite portions inside the austenite island are highlighted by circles. (b) Magnetization curve of DSS S32304 specimen (t = 0.69 mm, ε = -0.519) [Tavares et al. (2014)]. Further research on strain induced martensitic transformation was carried out on lean duplex grade S32304 and the metastability of austenite phase was studied by 38

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means of magnetic saturation (ms) measurements. Figure 2.24a shows the microstructure of DSS specimen of 1.01 mm thickness, undergone true cold deformation, ε = -0.138, with portions (circle) of αʹ-martensite inside the austenite islands. The ms value can be obtained by extrapolating the curve to H=0, as shown in Figure 2.24b. It has also been witnessed that the magnetization saturation increased with increasing cold deformation due to 𝛾 → 𝛼ʹ transformation [Tavares et al. (2014)]. Tavares et al. (2014) also reported that the lean duplex grades were much more prone to strain induced transformation than standard duplex grades. The austenite phases in these grades are even more metastable than austenitic grade AISI 304L. This is due to less Ni and Mo content in these steels. Several researchers concluded that plastic deformation had significant role to play in DSSs and it could influence intermetallic precipitation [Júnior & Balancin (2011), Choi et al. (2011), Farnoush et al. (2010)]. Some other group of researchers reported the deformation induced spinodal ferrite separation [Hättestrand et al. (2009)]. Deformation produces lattice defects and accelerates the diffusion within both phases which can act as potential sites for σ precipitation [Maehara et al. (1983)]. In nitrogen alloyed DSSs, the ferrite experiences more proneness towards decomposition under cold deformation [Weisbrodt-Reisch et al. (2006)]. Some authors reported the effect of solution annealing temperature on cold rolled structure. Figure 2.25 is showing the SEM micrographs of the DSS specimens aged at 800 °C for 1h after being solutiontreated at 1050 °C and 1250 °C [Cho & Lee (2013)].

Fig. 2.25: SEM micrographs of SAF 2205 aged at 800 °C for 1h as a function of solution treatment temperature; (a) 1050 °C and (b) 1250 °C. [Cho & Lee (2013)]

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From the micrographs, it could be pointed out that the morphologies of the σ phase were similar for both the solution treatment temperatures of 1050 °C and 1250 °C after successive plastic deformation and isothermal ageing. The precipitation of sigma is governed by higher driving force for precipitating intermetallics and by the higher diffusion rate of alloying elements after cold rolling [Cho & Lee (2013)]. Studies on annealing of cold rolled DSSs revealed that in case of ferrite, extensive recovery took place while annealing, whereas, austenite sub-structure remained almost same until the initiation of recrystallization. Despite of the higher driving force for recrystallization, the overall recrystallization kinetics was found to be less in austenite than in ferrite. Recrystallization nuclei are developed in ferrite by subgrain growth in a well recovered substructure. On the other hand, recrystallization in austenite takes place intermittently and preferentially at δ/γ interfaces [Reick et al. (1998)]. Recent developments on texture of cold rolling DSSs concluded the effect of microstructure prior to cold rolling on texture and formability. Previous findings suggest that the texture has strong dependency on the prior microstructure in (α+γ) micro-duplex stainless steel and a coarse α-grain structure promotes recrystallization with {111} initial orientation [Xiaoxu et al. (1995), Xiaoxu et al. (1996)]. Formation of shear bands in cold rolling was also reported when γ was present within α-matrix [Blicharski (1984)]. Recent investigation reveals that cold deformation results in strong α-fiber in δ-ferrite, which is a typical cold rolling texture of BCC metals. In γ-austenite, the strongest texture is MS-Brass with component {110} and this is also a typical cold rolling texture of FCC metals. This report supports the former investigation [Fargas et al. (2008), Keichel et al. (2003)]. The texture of the δ-phase in cold rolled and annealed DSSs were found to be affected by the structure prior to cold rolling as low temperature annealing (1050 °C compared to 1200 °C) produces more intense α-fiber. However, in case of γ-phase, the effect of initial microstructure was not clearly revealed. The microstructure prior to cold rolling has significant effect on the orientation of ferrite after final annealing. Specifically, components {100} and {211} decrease with coarsening of ferrite grains prior to cold rolling, whereas, microstructure prior to cold rolling does not play substantial role in final texture of austenite. The ODF of cold rolled and cold rolled annealed samples extracted from literature, are shown in Figure 2.26 [Hamada et al. (2010)]. It is witnessed from the ODFs that the grains having orientation || ND has been increased after annealing 40

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of δ-ferrite. In γ-austenite, two dominant hot rolling orientations || ND and || ND remain even after cold rolling and annealed sample [Hamada et al. (2010)]. Literature supports this finding [Xiaoxu et al. (1995)]. It was also reported that no recrystallization of the ferrite matrix took place in case of initial sub-grain size 25%). Besides, DSSs possess extremely good corrosion resistance property. Effectively, this 41

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combined effect of suitable mechanical property and corrosion resistance, DSSs became popular in recent years. High proof (Rp) and ultimate (Rm) strengths depend on several simultaneous mechanisms [Gunn (1997)]: • Interstitial (C, N) solid solution hardening • Substitutional (Cr, Mo, Ni, etc.) solid solution hardening • Strengthening by grain refinement due to the presence of dual phases • Possible hardening due to the formation of γ2 phase • Strengthening due to ferrite, since, for a similar composition, this phase is harder than the austenitic structure • Strain induced by differential contraction of the two phases on cooling from annealing temperatures The equations below were derived [Nordberg (1994)] from regression analysis. This held well with the experimental data [Charles (1994)]. 𝑅𝑝0.2 = 120 + 210√𝑁 + 0.02 + 2(𝑀𝑛 + 𝐶𝑟) + 14𝑀𝑜 + 10𝐶𝑢 +(6.15 − 0.054𝛿)𝛿 + {7 + 35(𝑁 + 0.02)}𝑑

−1⁄ 2

𝑅𝑚 = 470 + 600(𝑁 + 0.02) + 14𝑀𝑜 + 1.5𝛿 + 8𝑑

−1⁄ 2

(2.9) 𝑀𝑃𝑎

(2.10)

where, δ is the ferrite content in % and d is the lamellar spacing in mm. N has the most significant effect on tensile characteristics of DSSs. N is a austenite stabilizer and can preferentially strengthen austenite phase by interstitial solid solution strengthening, to the point where austenite becomes harder than ferrite [Wahlberg & Dunlop (1987)]. The tensile properties of DSSs are found to be influenced by plate thickness as anisotropy increased with reduction in plate thickness (Figure 2.27a) [Hutchinson et al. (1986)]. This anisotropy and strengthening associated with it are governed by refinement of the structure, with ‘duplex’ structure being fragmented parallel to the principal strain axis (Figure 2.27b). There is no such difference in hardness value between the phases, however, a marked unusual texture is observed with the major orientations of type (100)[011] to (211)[011] and (110)[223] for ferrite and austenite fibers respectively [Gunn (1997)]. It is seen that the variation in mechanical strength of DSSs is necessarily controlled by the ferrite. However, these alloys can undergo hardening of austenite due to strain induced martensitic transformation during tensile test [Cigada et al. (1991), Grundmann et al. (1988)]. Reports suggest that the cast DSS samples exhibit low strength owing to 42

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chemical segregation and presence of Widmanstätten austenite within δ matrix. 30% cold rolling has been found to increase the yield strength of the cast DSS about twice with only 10% total elongation and result into attractive combination of strength and ductility (UTS×%TEL = 25.22 GPa%) [Sugimoto (2000)].

Fig. 2.27: (a) Representative proof stress (Rp0.2) in the rolling and transverse directions, as function of plate thickness and (b) Petch plot relating proof stress (average of 0° and 90° tests) to lamellar spacing d. [Hutchinson et al. (1986)] The engineering stress-strain curves of DSSs subjected to different thermomechanical treatments are extracted from the literature and shown in Figure 2.28 [Ghosh et al. (2012)]. The solution annealed samples possess continuous yielding and higher uniform elongation, which is attributed to high austenite to ferrite ratio and globular shaped austenite after solution treatment.

Fig. 2.28: Engineering stress–strain curve of DSS samples after various treatments. [Ghosh et al. (2012)]

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Solutionizing may increase the ferrite content, but, refinement of the austenite may take place simultaneously. This is attributed to the fact the austenite is substantially stronger than ferrite. Residual stress may be present after annealing the cold rolled steels and hence they possess a bit higher yield stress value [Ghosh et al. (2012)]. Ferrite accommodates maximum strain at the early stages of tensile deformation [Balacin et al. (2000)]. But, at higher strain levels, load is shifted to austenite phase and led to dislocation density in this phase. A recent work emphasized more on the effect of temperature and strain rate on tensile property of lean duplex grade S32101 [Tsuchida et al. (2014)]. The nominal stress-strain curves obtained from literature are given in Figure 2.29.

Fig. 2.29: (a) Nominal stress–strain curves at various temperatures and (b) Nominal stress–strain curves with various strain rates at 296 K in the S32101 steels. [Tsuchida et al. (2014)] It can be noticed that the uniform elongation of S32101 substantially increase from 283 to 273 K, reaches peak value at 258 K, and decreases subsequently with decreasing temperature (Fig. 2.29a). Austenite phase undergoes martensitic transformation fractionally below 283 K and nearly all austenite islands transform to strain induced martensite below 173 K. Yield stress value is larger when high strain rate is applied in the process, whereas, uniform elongation is larger in strain rate below 10-2 s-1 at 296 K (Fig. 2.29b) [Tsuchida et al. (2014)]. Standard and lean DSSs can be hardened by short exposure at 475 °C, with a small expense of ductility and, therefore, heat treatments for 4, 8 or 12 hours are sometimes employed to increase the strength [Tavares et al. (2012)]. 44

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Literature Review

Fracture analysis To understand the micro-mechanism of fracture, tensile fracture surfaces have been studied by researchers. The fracture surface reveals dimples for ductile fracture and dimple size decreases with increasing ductility. The hot rolled annealed samples divulge finer dimples compared to the hot rolled samples and this features are in conformity with the percentage elongation values of samples. The cold rolled samples are strain hardened and, thus, reveals quasi-brittle fracture surface. This may be ascribed to the existence of arrested stress after heavy cold deformation which drives the motion of advancing crack. Cold rolled annealed samples reveal mainly ductile fracture with a range of dimple size and depth, which in turn, indicates recurrence of desirable ductile property [Ghosh et al. (2012)]. SEM micrographs of fracture surfaces of DSS samples subjected to different thermo-mechanical treatments are brought out from literature and shown in Figure 2.30.

Fig. 2.30: SEM micrographs of fracture surfaces in (a) hot rolled, (b) hot rolledannealed, (c) cold rolled and (d) cold rolled-annealed DSS samples. [Ghosh et al. (2012)]

45

Chapter 2

Literature Review

A paper published last year concludes that the warm-rolled DSSs possess ultrahigh strength owing to the presence of dislocations, ultra-fine structure with high fraction of HAGBs and phase boundaries. A sudden increase in ductility with expense of strength after annealing may be result of phase boundaries to be the only obstacle to dislocation glide [Zaid Ahmed & Bhattacharjee (2015)].

DSSs have an attractive combination of corrosion resistance and sufficient mechanical properties in the temperature range −50 to 250 °C, as they experience propensities towards embrittlement at temperatures above 250 °C [Hättestrand et al. (2009)]. Overall, these grades are suitable for use in many industries. An optimum alloy designing and a feasible processing route are essential prerequisites for extracting maximum benefit. Hence, a detailed study is performed on the effect of processing parameters on microstructure, microtexture and tensile properties of standard and lean DSSs in this work and is discussed elaborately in this thesis.

46

Chapter 3

EXPERIMENTAL DETAILS

3.1 Material 3.2 Microstructural evolution 3.3 Processing 3.4 Characterization

Chapter 3

Experimental Details

3.1 Material A standard DSS material was received from Shyam Ferro Alloys Ltd. in the form of 6 mm sheet in cold-rolled and annealed condition. Chemical composition of the asreceived standard DSS steel was determined using spectroscopic study and is given in Table 3.1. Before commencing the experimental studies, thermodynamic and kinetic simulations were performed using Thermo-Calc® (Version S) and JMatPro® (Version 7.0) to decide the composition of the cast material. A lean duplex grade material of around 2 kg was prepared in air induction furnace in the melting and casting laboratory at IIT Kharagpur. The chemical composition of the as-cast lean duplex grade is also presented in Table 3.1. Table 3.1: Chemical composition of the investigated duplex stainless steels (values are in weight %). Alloying elements

Sample Code

C

Si

Mn

Cr

Ni

Mo

N

Cu

Others

As-received steel (standard duplex grade)

SD

0.035

0.39

1.37

22.3

5.6

3.08

0.2

0.15

Co, V, W etc.

Laboratory melt (lean duplex grade)

LD

0.035

0.54

5.60

20.85

1.31

0.068

0.072

0.04

V, S (0.08) etc.

Sample source

While the as-received steel is a standard duplex grade with higher Ni content (5.5 weight %), laboratory made steel melt (lean duplex grade) is leaner with smaller amount of the costly alloying elements like Ni (1.31 weight %) and Mo. Sufficient S removal could not be done due to constraints in melting in air induction furnace. This is detrimental and may result in formation of MnS inclusion. The as-cast block of lean DSS grade (LD) was reheated to 1200 °C for 1 h inside a programmable muffle furnace, hot-forged down to ~ 17 mm thickness plate (applying ~ 66% reduction) at 1000-950 °C and finally air-cooled.

3.2 Microstructural evolution Cylindrical samples of 4 mm diameter and 10 mm height were prepared from standard DSS (SD) grade sheet and subjected to dilatometric study using BAHR®-

48

Chapter 3

Experimental Details

Thermoanalyse (DIL 805A/D model) dilatometer. Schematic diagram showing the thermal schedule for dilatometer tests is shown in Figure 3.1.

Fig. 3.1: Schematic diagram of the thermal schedule used for dilatometry. The samples were heated at two different heating rates (10/1 °C-s-1) to two different soaking temperatures (1000/900 °C) and cooled to room temperature at 5 °Cs-1.

3.3 Processing The processing schedules applied to different steels under investigation are shown schematically in Flow charts 3.1 and 3.2.

Flow chart 3.1: Processing route of the as-received material (standard duplex grade, SD). t is thickness.

49

Chapter 3

Experimental Details

Flow chart 3.2: Processing route of laboratory made as-cast steel (lean duplex grade, LD). t is thickness. 3.3.1 Heat treatment The as-received SD material was cut into small pieces and put into programmable muffle furnace to study the effect of isothermal holding temperature and time on the microstructure. In the first set of experiments, the samples were held for 30 minutes at different temperatures [1000 °C (IC1), 1050 °C (IC2), 1100 °C (IC3), 1150 °C (IC4), 1200 °C (IC5) and 1250 °C (IC6)]. In the second schedule, the samples were held at 900 °C for different times [0.25 h (IT1), 0.5 h (IT2), 1 h (IT3), 2 hrs (IT4), 5 hrs (IT5) and 8 hrs (IT6)]. The processing details of the heat treated SD samples, along with the sample codes are presented in Table 3.2. Upon 30 minute isothermal holding in the temperature range of 850-900 °C, sigma (σ) phase formation is expected in SD grade and not in LD grade. Besides that, 30 minutes holding time is typical in industrial 50

Chapter 3

Experimental Details

practice considering the sample size used in the present study. In a similar way, the nose of the TTT curve for the sigma (σ) phase lies around 900 °C for SD grade and thermodynamically 50-50 phase ratio is possible for LD grade at this temperature. All the samples were rapidly quenched in water to room temperature to freeze the high temperature structure. The schedule is schematically shown here in Figure 3.2.

Fig. 3.2: Schematic diagram of isochronal (IC) and isothermal (IT) heat treatment schedules. The as-forged LD material was also subjected to heat treatment. One sample was kept at 1200 °C for 30 minutes and cooled rapidly to room temperature to elucidate the high temperature structure and ferrite stability. Three samples were held at 900 °C for 1 h (LD_HT9001), 3 hours (LD_HT9003) and 5 hours (LD_HT9005) and one sample was held at 850 °C for 3 hours (LD_HT8503) and quenched directly in water to freeze the high temperature structure. The processing details of the heat treated LD samples, along with the sample codes are presented in Table 3.3. 3.3.2 Hot rolling The as-forged LD sample (t = 17 mm) was hot-rolled (i) down to 10.7 mm thickness (~37% reduction) using finish rolling temperature (FRT) of ~850 °C and (ii) down to 8.5 mm thickness (~50% reduction) using finish rolling temperature of ~520 °C. The soaking time of the LD steel samples at 1200 ˚C was 30 minutes, prior to hot rolling. The applied load was about 23-26 kN. FRT was recorded by digital pyrometer. The sample code of each rolled sample is mentioned in Table 3.3.

51

Chapter 3

Experimental Details

Table 3.2: Processing details of standard duplex (SD) steel samples. Sample code

Processing details

Isochronal treatment (IC) IC1

Sample soaked at 1000 °C for 30 minutes and then water quenched

IC2

Sample soaked at 1050 °C for 30 minutes and then water quenched

IC3

Sample soaked at 1100 °C for 30 minutes and then water quenched

IC4

Sample soaked at 1150 °C for 30 minutes and then water quenched

IC5

Sample soaked at 1200 °C for 30 minutes and then water quenched

IC6

Sample soaked at 1250 °C for 30 minutes and then water quenched Isothermal treatment (IT)

IT1

Sample soaked at 900 °C for 15 minutes (0.25 h) and then water quenched

IT2

Sample soaked at 900 °C for 30 minutes (0.5 h) and then water quenched

IT3

Sample soaked at 900 °C for 1 hour and then water quenched

IT4

Sample soaked at 900 °C for 2 hours and then water quenched

IT5

Sample soaked at 900 °C for 5 hours and then water quenched

IT6

Sample soaked at 900 °C for 8 hours and then water quenched

3.3.3 Cold rolling The 50% hot rolled sample LD_HR50 (8.5 mm) was subsequently cold-rolled down to ~ 4.8 mm thickness (~ 44% reduction) in a 4-hi cold rolling mill supplied by Tenova Multiform Ltd. at ambient temperature (25 ˚C) with a maximum applied rolling load of 75 kN. 3.3.4 Stress relieving annealing Industrially the cold rolled samples are annealed to relieve the stress, homogenize the structure and to restore the ductility of the material which gets deteriorated drastically by cold rolling. The hot rolled sample LD_HR50 and cold rolled sample LD_CR44 were subjected to annealing at 850 °C for different times (15 minutes and 3 hours). The processing details of rolled and annealed samples are presented in Table 3.3.

52

Chapter 3

Experimental Details

Table 3.3: Processing details of lean duplex (LD) steel samples. Sample code

Processing details

Heat-treated samples LD_HT1200

Sample soaked at 1200 °C for 30 minutes and then water quenched

LD_HT9001

Sample soaked at 900 °C for 1 hour and then water quenched

LD_HT9003

Sample soaked at 900 °C for 3 hours and then water quenched

LD_HT9005

Sample soaked at 900 °C for 5 hours and then water quenched

LD_HT8503

Sample soaked at 850 °C for 3 hours and then water quenched Rolled samples

LD_HR37

17 mm as-forged sample hot rolled by 37% down to 10.7 mm thickness, FRT ~850 °C

LD_HR50

17 mm as-forged sample hot rolled by 50% down to 8.5 mm thickness, FRT ~520 °C

LD_CR44

LD_HR50 sample cold rolled by 44% down to 4.76 mm thickness Rolled and heat-treated samples

HR_HT8500.25

Hot rolled sample LD_HR50 (t = 8.5 mm) heat treated at 850 °C for 15 minutes and then water quenched

HR_HT8503

Hot rolled sample LD_HR50 (t = 8.5 mm) heat treated at 850 °C for 3 h and then water quenched

CR_HT8500.25

Cold rolled sample LD_CR44 (t = 4.76 mm) heat treated at 850 °C for 15 minutes and then water quenched

CR_HT8503

Cold rolled sample LD_CR44 (t = 4.76 mm) heat treated at 850 °C for 3 h and then water quenched

3.4 Characterization After processing all the steel samples, they were characterized using various techniques; the details are given below. 3.4.1 Microstructural characterization Different microstructural characterization techniques were employed in order to elucidate the effect of processing parameters on microstructure of duplex stainless steels; these are listed below: 53

Chapter 3

Experimental Details

a. Sample preparation All the samples (dilatometer, heat treated, hot rolled and cold rolled) were mechanically ground, polished and etched with aqua regia solution (conc. HCl : conc. HNO3 3:1) and 10% aqueous oxalic acid solution (C2H2O4)

to observe the microstructure.

Microstructures were captured with Leica Optical Microscope (LEICA®DM6000M). In order to observe the σ phase, samples were etched by Villela reagent (chemical method). It contained 5 ml HCl, 1 gm Picric Acid (C6H3N3O7) and 100 ml Ethanol (C2H6O) (95%). Following Michalska & Sozańska (2006), etching was done for about 1:30 min at ambient temperature (25 ºC). The schematic diagram presented in Figure 3.3 shows the planes from where the samples were taken for different characterizations. b. Image analysis Image analysis was done and phase fraction, grain aspect ratio was estimated using software called Image J Analyzer. At least, six different frames with minimum 100 numbers of austenite islands were taken into consideration from each sample and the average value of phase fraction (%) with standard deviation is presented in this thesis. c. Hardness measurement The hardness was measured in Vickers Hardness Testing Machine. The applied loads were 25 gm-f/10 gm-f (HV0.025/HV0.01) and 1 kg-f (HV1.0) for microhardness and bulk hardness respectively and indentation time was 15 s. At least, twelve readings were taken for each sample and their average value with standard deviation is presented in this work. d. Lattice parameter determination For lattice parameter determination through X-ray diffraction (XRD) technique, Kα radiation from Copper target was used with a 2θ range of 40º - 130º. The step size was 0.02º and time per step was 0.5 second. The λ-value of the radiation is 1.67Å. For data analysis, direct comparison method was applied to find out the phases and an approximate estimation of their volume fractions. e. Chemical analysis Scanning electron microscope, SEM (Zeiss® EVO 60 model), fitted with Oxford-Inca ® PENTA FETX3 software for energy dispersive spectroscopic (EDS) analysis was used for microstructural characterization of the samples. Each phase was identified and chemical analysis of different phases was done separately. At least, eight points/areas 54

Chapter 3

Experimental Details

were taken from each phase of a sample to determine the chemical composition and their average value with standard deviation is presented in the thesis. Schematic diagram of DSS strip showing samples taken for different characterization techniques is given in Figure 3.3.

Fig. 3.3: Schematic diagram of DSS strip and samples taken for microstructural, EBSD and tensile property study. t is total thickness. 3.4.2 Microtextural characterization Samples were taken only from LD grade steel and microtextural characterization was carried out using EBSD technique. a. Grain boundary misorientation The sample preparation for EBSD consisted of standard mechanical polishing, followed by electro-polishing in 5% perchloric acid + 95% acetic acid solution. Step-size was 1.0 μm, covering 1000μm×750μm microstructural area on RD-TD plane, at the centre of mid-width location. EBSD (attached with Zeiss® EVO 60 SEM) characterized the phase map of δ-grains and γ-grains, grain boundary misorientation angle and local average misorientation. In the Orientation Image Mapping (OIM), the boundaries having misorientation angle of 15 degrees or more were considered as high-angle boundaries, i.e., ‘grain boundaries’. Boundaries having misorientation angles between 2 and 15 degrees were considered as low-angle boundaries, i.e., ‘sub-boundaries’.

55

Chapter 3

Experimental Details

b. Crystallographic orientation Orientation distribution function (ODF), inverse pole figure (IPF), twin boundary maps, boundaries with Kurdjumov-Sachs (K-S) orientation relationship, band contrast diagrams, strain contour maps etc. have been generated using EBSD. The analyses were carried out with softwares like HKL CHANNEL FIVE and TSL OIM Version 7.1. 3.4.3 Mechanical testing LD material was subjected to tensile test after different treatments. Plate-type sub-size (due to material constraint) tensile specimens were collected from heat treated sample LD_HT8503, hot rolled sample LD_HR50, cold rolled sample LD_CR44 and cold rolled annealed sample CR_HT8503. The specimen dimensions are shown in Figure 3.4. The basic consideration in designing the specimens was to maintain l/√A ratio greater than 4.0 following ASTM E-8M standard, where ‘l’ is the gauge length (GL) of the specimen, ‘A’ is the cross sectional area. All the tensile tests were conducted in Instron-8862 Universal Testing Machine (10 ton capacity) at room temperature (25 °C) using a cross-head velocity of 1 mm-minute-1. The GL of LD_HR50 tensile specimen was 25 mm and for other three specimens the GL was 20 mm due to shortage of material.

Fig. 3.4: Geometry of sub-size tensile specimens. Fractography and failure analysis Broken tensile specimens were seen under Scanning Electron Microscope. Both the Secondary (SE) and Back Scattered Electron (BSE) mode were used for fractography and failure analysis.

56

Chapter 3

Experimental Details

Microstructural and microtextural investigation of tensile tested specimens Two samples from the tensile tested specimens were taken for further texture investigation through EBSD to study the microtextural changes as well as microstructural changes during static tensile loading. The heat treated sample LD_HT8503 and the cold rolled annealed sample CR_HT8503 were taken into consideration. The samples were prepared from the reduced section area at a depth of about 1/4th of the specimen thickness on RD-TD plane and mechanically ground, polished and chemically etched in order to study the microstructural changes. Electro polishing was done to examine the final microtexture. The region of the tensile specimens from where the samples were taken for this study is shown schematically in Figure 3.5.

Fig. 3.5: Schematic diagram of the region (hashed zone) from where the tensile specimen was chosen for microstructural-microtextural study and fracture analysis. t is total thickness.

57

Chapter 3

Experimental Details

58

Chapter 4

PHASE DIAGRAM AND MICROSTRUCTURAL EVOLUTION

4.1 Introduction 4.2 Calculation of Phase transformation 4.3 Microstructural characterization 4.4 Microstructural evolution 4.5 Summary & Conclusion

Chapter 4

Phase Diagram and Microstructural Evolution

4.1 Introduction In order to examine the microstructural evolution in duplex stainless steel as a function of heat treatment schedule, phase transformation calculations have been carried out using Thermo-Calc® (Version S) and JMatPro® (Version 7.0). Depending upon the result of simulation study, further processing parameters were chosen for the heat treatment trials of composition SD and LD (as mentioned in section 3.1). To study the microstructural evolution, two simple dilatometry cycles have also been applied, as mentioned in section 3.2. Theoretical calculations depending upon the lattice parameters of the phases were conducted to justify the result obtained from the dilatation curves.

4.2 Calculation of Phase transformation Duplex stainless steels are basically Fe-Cr-Ni alloys and to understand the evolution of different phases with varying temperature thermodynamic and kinetic calculations need to be carried out. In order to do this, Fe-Cr-Ni pseudo ternary diagrams at different temperatures and phase fraction diagrams of the examined steels have been plotted by thermodynamic software Thermo-Calc® and TTT, CCT diagrams have been constructed using JMatPro®. The microstructure has been predicted by Schneider diagram after calculating the Chromium equivalent (Creq) and Nickel equivalent (Nieq). 4.2.1 Fe-Cr-Ni pseudo-ternary diagrams Schematic diagram of Fe-Cr-Ni system with 70% Fe level was constructed to identify the phase proportions and their compositions [Schafmeister & Ergang (1939)]. In this study, an approximate composition of 70Fe-23Cr-6.5Ni was chosen and pseudo ternary diagrams have been plotted at different temperatures with three axes denoting the mole fractions of Fe, Cr and Ni. These are basically two-dimensional sections extracted from the complex three-dimensional ternary diagram over the whole temperature range. The calculated Fe-Cr-Ni pseudo ternary diagrams at 1200, 1100, 1000, 900, 800 and 650 °C are shown in Figure 4.1.

60

Chapter 4

Phase Diagram and Microstructural Evolution

Fig. 4.1: Fe-Cr-Ni pseudo ternary diagrams at (a) 1200 °C, (b) 1100 °C, (c) 1000 °C, (d) 900 °C, (e) 800 °C and (f) 650 °C. (Computed using Thermo-Calc®) The black dot indicates the composition of the materials used in this study. 61

Chapter 4

Phase Diagram and Microstructural Evolution

The composition of the standard DSS materials used in this study is nearly 70Fe23Cr-6.5Ni weight % and is shown by black dot inside the diagram. At 1200 °C, the point lies approximately along the α-solvus line. As the temperature went down, the equilibrium amount of austenite phase is expected to increase and upon very long holding at that particular temperature, the equilibrium amount of the phases can be obtained. The point of consideration corresponding to the present SD alloy composition shifted from α-solvus line to the two-phase region (ferrite + austenite) with decreasing temperature. The point tended to shift from two-phase to three-phase region as the temperature decreased and this third phase was deleterious sigma phase. Around 800 °C, this composition started to show the formation of intermetallics (sigma). At 650 °C, the microstructure would be austenite in ferrite matrix and a small amount of sigma phase would also be present. The amount of sigma phase depends on the amount of Cr and Mo added to the steel as these are the main elements responsible for formation of sigma phase. Further decrease in temperature resulted in formation of other intermetallic and undesired phases [Solomon & Devine (1979), Williams (1958)]. The amount of intermetallic phases depends on the amount of alloying elements and their partitioning coefficient [Charles (1994), Hertzman et al. (1992)].

4.2.2 Phase fraction diagrams The presence of alloying elements other than Cr and Ni makes substantial modification in phase fields. Taking these elements in consideration, one approach has been to make by pseudo-ternary sections of multi-component alloy (for example, Fe-Cr-Ni-Mo-N) which is used to calculate the phase equilibria over a wide temperature range [Nilsson (1992)]. This could then be used to form an isopleth diagram for a given composition and this approach has been found to be in good agreement with experimental data [Gunn (1997)]. The phase fraction diagrams indicate the amount of thermodynamically stable phases in a temperature range and these have been drawn using Thermo-Calc® for the steels examined in this work. The phase fraction diagrams of the investigated steels are shown in Figure 4.2. It indicated that the amount of ferrite increased as the temperature went up from around 900 °C. The main focus of the study was to achieve more or less 50% of each phase. The red line indicated the temperature in which more or less equal proportion of δ-ferrite and γ-austenite was stable under equilibrium condition. This temperature was around 1080 °C and 900 °C for steels SD and LD,

62

Chapter 4

Phase Diagram and Microstructural Evolution

respectively. The phase fraction data at 50 °C interval for LD material are presented in Table 4.1.

Fig. 4.2: Phase fraction diagrams of composition (a) SD and (b) LD. (Computed using Thermo-Calc®) Table 4.1: Phase fraction data calculated by Thermo-Calc® for steel LD at 50 ˚C interval (values are in volume %). Temperature (ºC)

γ

δ

σ

Cr23C6

(Cr,V)2N

Cu-rich ε

100

1300 1250

0.5034

99.497

1200

9.3722

90.628

1150

18.518

81.482

1100

27.549

72.451

1050

35.647

64.34

0.012506

1000

40.852

59.05

0.09863

950

45.436

54.39

0.17391

900

49.138

50.623

0.23898

850

50.662

48.883

0.16122

0.29387

800

50.519

48.778

0.36397

0.33822

750

57.744

32.037

9.3623

0.48298

0.37404

700

72.257

26.783

0.55291

0.40702

650

53.154

17.744

28.039

0.62914

0.43316

600

20.993

45.608

32.279

0.66731

0.45293

550

4.1467

59.245

35.473

0.677

0.45833

500

0.24859

62.852

35.758

0.68098

0.45997

450

63.504

35.334

0.68361

0.46179

0.016429

400

63.292

35.534

0.68513

0.46369

0.025539

350

62.57

36.248

0.68609

0.4656

0.030452

300

61.633

37.179

0.68679

0.4675

0.032927

63

Chapter 4

Phase Diagram and Microstructural Evolution

Likewise, the phase fraction diagrams of composition SD was analyzed at different temperatures. 4.2.3 Time-Temperature-Transformation (TTT) curves Time-Temperature-Transformation (TTT) diagrams, produced by isothermal heat treatment followed by quenching, are often employed to depict the susceptibility of different grades of stainless steel to embrittlement. For example, Charles (1991) determined the TTT diagram of standard grade duplex stainless steel. The TTT curve was based on optical microscope observations in the temperature range of 600-1050°C and on hardness measurements in the temperature range of 300-600°C [Charles (1991)]. The duplex alloy experiences different intermetallics and precipitation of deleterious phases in different temperature regions [Solomon & Devine (1979)]. Hence, to study the phase change behavior, TTT diagrams need to be constructed first. The kinetic calculations were done for all the investigated steel compositions using JMatPro® and the resultant TTT diagrams are shown in Figure 4.3.

Fig. 4.3: TTT diagrams of (a) SD and (b) LD steel grades, computed by JMatPro®. The as-received steel SD was richer in Mo as compared with the laboratory made lean duplex grade LD. This clearly showed why SD grade was more prone to sigma phase precipitation. The nose of the sigma start curve was in the temperature range of 850-900 °C and at 570 °C for composition SD and LD, respectively. The sigma (σ) phase would precipitate within 2 minutes at the nose temperature for SD grade steel as the rate of diffusion would be quite high at that temperature. In case of 64

Chapter 4

Phase Diagram and Microstructural Evolution

LD composition, this would take as long as 4 days at the nose temperature due to poor rate of diffusion. The precipitation behavior of σ phase is determined by both the concentrations of σ forming elements, such as Cr and Mo, in the ferrite and the δ/γ ratio before precipitation [Maehara et al. (1983)]. Like σ-phase, χ-phase forms between 700 and 900ºC, although in much smaller quantities [Nilsson (1992)]. However, enrichment of ferrite with intermetallic forming elements during a long exposure to relatively low temperatures, i.e. 700°C, favors the precipitation of χ-phase [Gunn (1997)]. In case of the investigated steel SD grade, the minimum calculated time to precipitate χ-phase was 36 seconds at the nose temperature of about 850 °C. Cr2N is more likely to form after higher solution heat treatment temperatures [Roscoe et al. (1984)] and forms rapidly even if quenched from such temperatures [Nilsson & Liu (1988)]. Isothermal exposure at 700-950 °C temperature range produces intergranular Cr2N at δ/δ grain boundaries. In case of investigated lean duplex grade LD, M2(C,N) was predicted to precipitate after 20 minutes of holding at 800 °C. In duplex grades with moderately high carbon levels of about 0.03%, the carbide M23C6 rapidly precipitates between 650 and 950 °C, requiring less than one minute to form at 800 °C [Solomon (1982)] for alloy S32304. In steel LD, the TTT curve shows that the carbide particles would start to precipitate after about 80 minutes holding at 750 °C. Short exposure in this temperature range might not result in carbide precipitation. The lowest temperature decomposition product in duplex steel is that of α′, which occurs between 300 and 525°C, and is the main cause of hardening and ‘475 °C embrittlement’ [Fisher et al. (1953), Blackburn & Nutting (1964)]. It is suggested that α′-formation is a consequence of the miscibility gap in the Fe-Cr system [Williams (1958)], whereby ferrite undergoes spinodal decomposition [Lagneborg (1967)]. The TTT diagram of steel LD clearly shows that the start of formation of α' phase takes about 25 hours at 420 °C. 4.2.4 Continuous Cooling Transformation (CCT) curves Isothermal cooling is not practiced in industry and all the steel components (e.g. slab, ingot) are cooled continuously from high temperature to room temperature. Hence, it is important to investigate the results pertaining to continuous cooling situation. At

65

Chapter 4

Phase Diagram and Microstructural Evolution

temperatures near the solvus, the nucleation of precipitates is slow and their growth is fast, whereas the opposite is true at lower temperatures, near the ‘nose’ of the transformation curve. Therefore, it is difficult to avoid phase transformations, such as sigma precipitation, during the reheating of heavy section products (e.g. ingots, castings, thick plate, etc.), and so a solution treatment should be performed at a sufficiently high temperature to redissolve any such phases. On the other hand, during cooling, the slow nucleation rate at high temperature and the sluggish growth rate at lower temperatures make it relatively easy to avoid the formation of sigma phase, even in the case of air cooling of certain castings or heavy plate [Gunn (1997)]. A continuous cooling transformation (CCT) diagram of the super duplex grade S32520 clearly shows a considerable shift in the position of the precipitation ‘nose’ between the TTT and CCT diagrams [Charles (1991), Josefsson et al. (1991)]. Cooling rates as low as about 2000 °C-h-1 (0.56 C-s-1) are possible without precipitation when solution annealing is done at 1080 °C. The CCT diagrams of investigated compositions are presented in Figure 4.4.

Fig. 4.4: CCT diagrams of (a) SD and (b) LD steel grades, computed using JMatPro®. Cooling curves for four different cooling rates, i.e., 0.028, 0.28, 2.8 and 28 C-s-1 are presented. In case of composition SD, if the solution annealing would be done at around 1000 °C and the steel would be allowed to cool from that temperature, sigma phase would start to precipitate at around 900 °C with a slow cooling rate of 0.28 C-s-1. Deleterious χ-phase would be kinetically more favorable than σ-phase to precipitate 66

Chapter 4

Phase Diagram and Microstructural Evolution

even when a rapid cooling rate as fast as 2.8 C-s-1 (faster than air cooling) is employed. But, since χ-phase forms in tiny amount, this has lesser effect than σ-phase as far as material property deterioration is concerned. Lean duplex LD grade is expected to experience lower propensities to intermetallic precipitation as this grade was leaner with alloying elements like Mo and Ni which are mainly responsible for formation of brittle, deleterious intermetallics. The level of N was also very less compared to SD grade. According to the calculated CCT curve of M2(C,N) it is expected to form at around 400 °C at a slow cooling rate of 0.28 C-s-1. However, the rate of diffusion would be very slow below this temperature so that negligibly small amount of precipitate is expected, if it forms at all. Similarly, lower Cr and C concentration of LD material would result in sluggish formation of M 23C6 carbides. The calculated CCT curve shows that M23C6 would start forming at 700 °C at as slow a cooling rate as 0.028 C-s-1. Isothermal heat treatment at 865 °C for 134 seconds precipitates 1% σ, which is also obtained at room temperature after continuous cooling at 0.23 °C-s-1 from a temperature above 930 °C in 21.5Cr-5.5Ni-3Mo-1.5Mn-0.17N steel [Sieurin & Sandström (2007)]. Due to leaner chemistry, the composition LD would not experience σ precipitation even when it is allowed to cool very slowly inside the furnace. DSSs experience precipitation of α' phase between 300 and 525°C, which is the main reason for ‘475 °C embrittlement’ [Blackburn & Nutting (1964)]. Investigated composition LD would start experiencing the formation of α′ phase at around 480 °C, but at that temperature, the available free energy for transformation would be very less. Extreme slow cooling rate (15° as high angle grain boundaries. The peak relative frequency value was 0.4 for 2° misorientation of both the phase boundaries, i.e., most of the boundaries were low angle grain boundaries (LAGBs). These LAGBs were formed owing to the recovery of dislocation formed during hot rolling [Dehghan-Manshadi et al. (2007)]. As the hot rolled samples were subjected to annealing treatment, the LAGBs started to move in order to relieve the strain accumulated during hot deformation. In case of δ-ferrite, the angle of misorientation increased gradually as the annealing time increased.

Annealing produced newly formed δ-grains with high angle grain

boundaries (HAGBs). The average angle of misorientation of δ grain boundaries was the highest at around 32° in sample HR_HT8503. These were essentially HAGBs. By contrast, austenite grain boundaries did not change and remained LAGBs up to a certain period of holding at 850 °C. After a prolonged heat treatment, a peak at 60° angle of misorientation was observed in sample HR_HT8503. This was a characteristic peak that came from annealing twin generated within γ-grains during isothermal holding. Hence, the microstructure evolution during annealing of hot deformed structure was very different in each phase. Continuous dynamic recrystallization (continuous DRX), i.e., gradual increase in angle of misorientation and distance between low angle grain boundaries was evidenced in δ-ferrite, whereas, newly formed recrystallized γ-grains were observed on deformed twin boundaries or δ/γ interphase boundaries, formed through conventional dynamic recrystallization (conventional DRX). Nonetheless, owing to limitation in the number of γ/γ grain boundaries, the extension of DRX in austenite was restricted. These findings agreed with the work of Dehghan-Manshadi et al. (2007). Local average misorientation The local average misorientation has been calculated to elucidate the strain accumulated by a grain locally with respect to the neighboring grains in terms of angle of misorientation. The relative frequency of occurrence of local average misorientation has been plotted in Figure 5.15. It was clearly evidenced that the hot rolled sample LD_HR50 had most number of δ-ferrite grains with about 1° angle of local misorientation, whereas, in case of γ-austenite phase, this value was 2°. It suggested 100

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that γ-austenite has been more distorted than δ-ferrite. But, as the annealing progressed at 850 °C, the relative frequency of occurrence rapidly increased at a very low angle in ferrite, which, in turn suggested quick strain release in this phase. In other words, ferrite phase rapidly resulted in complete softening due to the recrystallization phenomenon. Previous results [Belyakov et al. (2005)] recommended that the structural changes taking place in ferrite phase at the beginning of the continuous recrystallization (similar to sample HR_HT8500.25 of the present study) were qualitatively similar to those occurring at the nucleation stage of the primary recrystallization. Specifically, the fast recovery in the ribbon-like deformation substructures at an early annealing resulted in the rearrangement and annihilation of dislocations within the strain-induced continuously curved crystallites, leading to the breakup of the highly elongated subgrains into the chains of almost equiaxed sub-grains. The local misorientation maps of BCC phase in hot rolled and hot rolled annealed samples are given in Figure 5.16.

Fig. 5.15: Graphical representation of local average misorientation angle’s relative frequency of occurrence in sample LD_HR50, HR_HT8500.25 (850 °C, 15 minutes) and HR_HT8503 (850 °C, 3 h) – (a) δ-ferrite and (b) γ-austenite. The color code represents the grain in terms of misorientation angle in local misorientation map.

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Fig. 5.16: The local misorientation maps of δ-ferrite phases in hot rolled sample (a) LD_HR50 and hot rolled annealed samples: (b) HR_HT8500.25 (850 °C, 15 minutes) and (c) HR_HT8503 (850 °C, 3 h). The color code bar indicates the grain of particular angle of misorientation (in degrees). Rolling direction is parallel to the marker. In LD_HR50, ferrite phase consisted of grains with high angle of misorientation mostly on δ/γ interphase boundaries. The grains which were locally surrounded by ferrite neighbors underwent less amount of deformation and hence less distorted lattices were developed in these regions during hot rolling. As annealing took place, the ferrite phase readily got strain free. Ferrite regions within the closely fragmented austenite islands remained strained after initial stage of annealing in sample HR_HT8500.25. This was because the austenite phase got more strain hardened, i.e., more misoriented locally, the ferrite regions nearby this could not relieve the stress in their surroundings and stress remained there. Nevertheless, prolonged annealing treatment resulted in mostly strain free ferrite lattice. Since, 𝛿 → 𝛾 transformation took place while annealing, the newly formed γ enveloped the remaining δ phase. These δ/γ interfaces within large γ grains remained strained even after prolonged annealing treatment in sample HR_HT8503. However, in case of austenite phase, a gradual decrease in peak angle of misorientation has been witnessed. From earlier work by Belyakov et al. (2002), it has been understood that the average sub-grain boundary misorientation angle might remain unchanged after a certain time of annealing, the internal stress (𝜏⁄𝐺 ) decreased gradually as soon as the deformed austenitic structure (FCC) was subjected to isothermal annealing treatment. The local misorientation maps of austenite phase in hot 102

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rolled and hot rolled annealed samples are provided in Figure 5.17. These maps also revealed the same phenomena. The hot rolled sample LD_HR50 consisted of more number of deformed grains and the peak relative frequency was at around 2° (green). The deformity of the grains started to decrease as the isothermal annealing initiated. Eventually, the peak relative frequency of misorientation angle shifted to lower value. After prolonged annealing treatment (HR_HT8503), the newly formed tiny and recrystallized γ-grains appeared strain-free, whereas, some residual stress still remained inside the large grains. This might be the result of presence of internal stress within the lattice as it decreased rapidly up to a certain period of time and then the residual stress remained even after a long exposure at that temperature [Belyakov et al. (2002)]. This finding supported the presence of strained region in austenite phase.

Fig. 5.17: The local misorientation maps of austenite phases in hot rolled sample (a) LD_HR50 and hot rolled annealed samples: (b) HR_HT8500.25 (850 °C, 15 minutes) and (c) HR_HT8503 (850 °C, 3 hours). The color code bar indicates the grain of particular angle of misorientation (in degrees). Rolling direction is parallel to the marker. 103

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5.4 Cold rolled and cold rolled-annealed steels The effect of cold rolling and subsequent annealing treatment on microstructure of lean DSS grade steel (LD) was investigated. 5.4.1 Cold deformation The hot rolled sample LD_HR50 (t=9mm) was cold rolled down to a thickness of 5 mm at ambient temperature with a total thickness reduction of 44%. The light optical microstructure of cold rolled sample LD_CR44 is given in Figure 5.18a and image analysis results are shown in Table 5.7.

Fig. 5.18: Optical microstructures of cold rolled sample (a) LD_CR44 and cold rolled annealed samples: (b) CR_HT8500.25 and (c) CR_HT8503 (annealed at 850 °C for 15 minutes and 3 hours respectively) and (d) closer view of sample CR_HT8503. Rolling direction is parallel to the marker.

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Table 5.7: Results obtained from optical image analysis and hardness measurement of cold rolled sample LD_CR44. Volume % Sample ID

Avg. aspect ratio,

Phase γ-austenite

Longitudinal

Transverse

𝑨𝑹 = 𝒅𝒎𝒂𝒙 ⁄𝒅𝒎𝒊𝒏

23.5 ± 1.33

20.3 ± 1.13

8.33 ± 0.04

Microhardness (Hv) 466 ± 6.15

Avg. bulk hardness (Hv)

430 ± 4.78

LD_CR44 δ-ferrite

76.5 ± 1.33

79.7 ± 1.13

-

417 ± 6.53

Result of image analysis showed that the volume fraction of each phase remained almost the same as the hot rolled sample LD_HR50 as phase transformation at this temperature was not feasible due to less available energy for diffusion.

Deformation induced martensite Only (α + γ) ‘duplex’ microstructure was observed by light optical microscopy. However, previous work suggested that diffusionless transformation is possible while cold deformation of lean duplex stainless steel UNS S32304 [Tavares et al. (2014)]. Extensive study on deformation induced martensitic transformation of austenite (𝛾 → 𝜖 + 𝛼′ → 𝛼′) was carried out by Tavares et al. (2014). Since, ε-martensite is metastable compared to αʹ martensite, the final microstructure exhibited only the latter form. Same type of lath martensite has been witnessed within austenite islands in cold rolled sample LD_CR44 in secondary electron imaging. Taking the low carbon concentration of austenite into account, it could be said that the strain induced martensite was expected to form in lath shape. The secondary electron image of LD_CR44 is provided in Figure 5.19a, which clearly revealed the laths of αʹ-martensite within austenite grains. The measured volume % of austenite (islands) reported in Table 5.7 is therefore an overestimate since strain/deformation induced martensitic transformation occurred inside the islands during cold deformation and the volume % of αʹ-martensite needed to be measured separately. A relationship between the extent of thickness reduction (%) by cold rolling and the volume % of deformation induced αʹ-martensite was reported [Baldo & Mészáros (2010)] in a composition of DSS similar to LD. Hence, a rough estimation could be made from the graph furnished in that report. The plot is shown in Figure 5.21b; the amount of αʹ-martensite expected to form in the LD material after 44% thickness reduction was estimated to be around 8%. Since this α′-martensite forms at the expense of austenite, the actual amount of austenite remaining untransformed in cold rolled material was 15.57%. The volume % of δ-ferrite remained unchanged. 105

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Further image analysis was done by EBSD technique to find the exact volume % of ferrite and austenite. The result is discussed in next section in detail.

Fig. 5.19: (a) Secondary electron image of cold rolled sample LD_CR44, lath of αʹmartensite are visible inside γ-islands and (b) αʹ-martensite (volume %) versus thickness reduction (%) curve [Baldo & Mészáros (2010)]; amount of αʹ-martensite is shown after 44% thickness reduction. 5.4.2 Effect of annealing on cold rolled microstructure Two samples from the cold rolled sheet LD_CR44 were annealed isothermally at 850 °C for 15 minutes and 3 hours and quenched in water (mentioned in section 3.2.4). The microstructures of the annealed samples CR_HT8500.25 and CR_HT8503 are shown in Figure 5.18b and 5.18c respectively. The austenite islands appeared elongated but some 106

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islands tended to take globular shape (aspect ratio decreased). Prolonged annealing treatment resulted in some tiny austenite grains in sample CR_HT8503. A closer view of this sample is shown in Figure 5.18d, where newly formed γ grains could be seen. These petite γ grains were not observed in sample CR_HT8500.25. The results of image analyses of light optical images are given in Table 5.8.

Table 5.8: Results of image analyses obtained from optical microstructures of cold rolled and annealed samples CR_HT8500.25 (850 °C, 15 minutes) and HR_HT8503 (850 °C, 3 hours). Volume % Sample ID

Phase

Avg. aspect ratio,

Longitudinal

Transverse

𝑨𝑹 = 𝒅𝒎𝒂𝒙 ⁄𝒅𝒎𝒊𝒏

γ-austenite

29.3 ± 1.22

28.4 ± 0.89

2.86 ± 0.13

δ-ferrite

70.7 ± 1.22

71.6 ± 0.89

-

γ-austenite

32.5 ± 1.12

31.3 ± 1.27

1.96 ± 0.11

δ-ferrite

67.5 ± 1.12

68.7 ± 1.27

-

CR_HT8500.25

CR_HT8503

The phase fraction of austenite increased upon annealing. Since the equilibrium volume fraction of austenite at 850 °C is higher than that in the cold rolled sample, α′martensite (formed due to strain induced transformation during cold rolling) and part of δ-ferrite transformed to austenite. The change in hardness and average aspect ratio of cold rolled material with annealing time at 850 °C is shown graphically in Figure 5.20. The graph revealed that the microhardness of each phase as well as the bulk hardness of the steel gradually decreased with annealing time at 850 °C. The smooth boundary of austenite islands became wavy after long thermal exposure and could be seen in high resolution optical microstructure of sample CR_HT8503 (Figure 5.18d). Preferential growth of austenite from δ/γ interphase boundary towards ferrite grains might result in this type of microstructure. The nucleation and growth of newly formed austenite occurs by rejecting excess Cr and Mo (ferrite formers) to the neighborhood and accepting Ni and Mn (austenite formers) from the adjacent regions. This mechanism of nucleation and growth might explain the wavy nature of the interphase boundary. The average aspect ratio (AR) was also inversely proportional to the annealing time. 107

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Fig. 5.20: Comparison of the bulk hardness, microhardness of each phase and average aspect ratio of γ-austenite of the annealed samples CR_HT8500.25 (850 °C, 15 minutes) and CR_HT8503 (850 °C, 3 hours) with the cold rolled sample LD_CR44. Since the estimation of phase volume % was just an approximation by light optical microscopy and newly formed small γ-grains were not discovered efficiently, image analysis by EBSD technique was carried out. Like hot rolled samples, a comparative analysis was performed between the cold rolled and cold rolled + annealed materials in terms of grain boundary misorientation and local average misorientation. Phase map The phase maps obtained from cold rolled and cold rolled + annealed samples by EBSD technique are shown in Figure 5.21. The cold rolled sample LD_CR44 clearly showed deformed structure as dislocations piled up against grain boundaries. Since, volume % of austenite was overestimated by optical microscopy due to strain induced transformation of αʹ-martensite within austenite islands, phase volume % of sample LD_CR44 was interpreted again from the phase map. It revealed that the cold rolled sample consisted of 85.5 volume % ferrite phase and 14.5 volume % austenite phase. The result clearly conveyed that some portion of austenite (FCC) transformed to αʹmartensite (BCC) during cold rolling. To be specific, 9 volume % of austenite was 108

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found to transform to α′-martensite after 44 % thickness reduction. In previous work, the transformed volume % of austenite to αʹ-martensite was reported as 7.93% (mentioned in section 5.4.1) [Baldo & Mészáros (2010)] after 44% thickness reduction. But that study was based on single pass rolling and multi-pass cold rolling was carried out in this work. The applied load would be higher in case of multi-pass cold deformation and this resulted in slightly higher volume % of strain induced αʹmartensite than the estimated value obtained from the plot described by Baldo & Mészáros (2010). The annealed samples CR_HT8500.25 and CR_HT8503 revealed small sized ferrite grains and newly formed austenite grains along δ/δ grain boundaries that grew towards ferrite grains by consuming it. As the equilibrium volume % of austenite at 850 °C was higher than the volume % of austenite in cold rolled sample LD_CR44, δ to γ transformation took place. The strain induced martensite also transformed back to austenite when the sample exposed to elevated temperature. The average grain size of ferrite decreased and number of small austenite grains increased with increasing annealing time at 850 °C. The measurement of austenite grain average aspect ratio was also confirmed from the phase maps of all three samples.

Fig. 5.21: Phase maps of cold rolled sample (a) LD_CR44 and cold rolled + annealed samples: (b) CR_HT8500.25 (850 °C, 15 minutes) and (c) CR_HT8503 (850 °C, 3 h), obtained from EBSD. Rolling direction is parallel to the marker. 109

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To study the phase transformation mechanism during annealing the cold rolled sample and to establish a comparative analysis between cold rolled and cold rolled annealed samples, grain boundary misorientation and local average misorientation has been plotted and discussed in next subsections. Grain boundary misorientation The average grain boundary misorientation of individual phase has been measured and is plotted against relative frequency of occurrence in Figure 5.22. The plot revealed that the cold rolled sample LD_CR44 consisted of heavily deformed γ and δ grains with very low angle grain boundaries. Grain boundaries with 2-15° and >15° angle of misorientation were considered as low angle and high angle grain boundaries respectively. The peak relative frequency value was 0.25 for 2° misorientation of both the phase boundaries, i.e., most of the boundaries were low angle grain boundaries (LAGBs). This result held well with the previously reported findings [Maki et al. (2001)]. In case of austenite phase, some portions of austenite grains transformed to αʹmartensite due to application of stress during cold rolling. Martensite exhibited highly distorted lattice and thus, mostly surrounded by LAGBs. A small peak of relative frequency value 0.1 was observed at 60° in cold rolled sample LD_CR44, which could be a characteristic peak of deformation twin in heavily deformed austenite phase.

Fig. 5.22: The comparative grain boundary misorientation graphs of LD_CR44, CR_HT8500.25 (850 °C, 15 minutes) and CR_HT8503 (850 °C, 3 hours) in terms of relative frequency of occurrence – (a) ferrite and (b) austenite phase. Previous research work has suggested that the major strain between δ-ferrite (BCC) and γ-austenite (FCC) is accommodated in those areas where phase boundaries 110

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separate ferrite and austenite in rolling direction [Keichel et al. (2003)]. Practically, these areas amount to only a small fraction of the total interface, whereas, the largest part of the total phase boundary area is parallel to the sheet plane and as long as the strains in both phases were similar, no strong interaction between them was witnessed. As the cold rolled samples were subjected to annealing treatment, the thermal energy activated LAGBs to advance in order to alleviate the stress gathered during cold deformation. In case of δ-ferrite, the propensity towards recovery was faster than γaustenite due to the continuity in this mechanism. As an effect, annealing produced recovered δ sub-grains with high angle grain boundaries (HAGBs) just after 15 minutes of annealing at 850 °C. The average angle of misorientation of δ grain boundaries was the highest at around 45° in sample CR_HT8503. These were essentially HAGBs. It was seen that the recrystallization of δ sub-grain boundaries was inhibited by fine γ particles along the δ/δ interfaces and that resulted in a very stable structure by the long time aging treatment [Maki et al. (2001)]. The recrystallization of δ sub-grain matrix hardly occurred even during annealing after heavy cold rolling, as reported by Maki et al. (2001). Similarly, the main part of the newly formed γ-grains (FCC) during annealing was equiaxial and exhibited a low dislocation density. The presence of annealing twins was the main feature in case of austenite grains [Fargas et al. (2008)]. In contrast to ferrite, the microstructure development in austenite during annealing was evidently due to primary recrystallization [Keichel et al. (2003)]. After a prolonged heat treatment, the austenite grains revealed more grains with 60° angle of misorientation in sample CR_HT8503 because annealing twins were formed within γ-grains during isothermal holding. Hence, the microstructure evolution during annealing of cold deformed structure was very different in each phase and this phenomenon seconded the mechanism of microstructural evolution in hot rolled and hot rolled annealed materials. However, unlike hot rolled-annealed samples, gradual increase in angle of misorientation and distance between low angle grain boundaries was not evidenced in δ-ferrite. Recovery prevailed for a long time of annealing and delayed the onset of dynamic recrystallization (DRX), whereas, newly formed recrystallized γ-grains were observed on deformed twin boundaries or δ/γ interphase boundaries, formed through conventional dynamic recrystallization (conventional DRX). As the accumulated strain during cold rolling was much higher than in the hot rolled sample, the γ-phase experienced an inclination to faster nucleation and growth, which resulted in faster formation of HAGBs after only 15 minutes of annealing at 850 °C. 111

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Local average misorientation The local average misorientation has been estimated to elucidate the strain accumulated by a grain locally with respect to the neighboring grains in terms of angle of misorientation. The relative frequency of occurrence of local average misorientation has been plotted in Figure 5.23.

Fig. 5.23: Graphical representation of relative frequency of occurrence of local average misorientation angle in samples LD_CR44, CR_HT8500.25 (850 °C, 15 minutes) and CR_HT8503 (850 °C, 3 hours) – (a) ferrite and (b) austenite phase. The color code represents the grain in terms of misorientation angle in local misorientation map. It was clearly spotted from the graph provided in Figure 5.25 that in case of ferrite phase, the cold rolled sample LD_CR44 had the highest peak of relative frequency at about 2.7° angle of misorientation, whereas, peak of austenite phase came at near 2.3°. It suggested that austenite phase, which possessed higher hardness, has been less distorted than BCC (ferrite). In all cases, the average value of angle of local misorientation was higher in cold rolled steel than in hot rolled steel since the cold deformation was done at ambient temperature with higher rolling load. As the annealing initiated at 850 °C, the relative frequency of occurrence rapidly increased at a very low angle in BCC, which, in turn suggested quick strain release in this phase. It has already been discovered that the overall recrystallization kinetics was faster in ferrite than in austenite, in spite of the higher driving force for recrystallization in austenite [Reick et al. (1998)]. The recovered (partially recrystallized) structure showed a peak at very low angle of misorientation after prolonged annealing in sample CR_HT8503. The local misorientation maps of ferrite phase in cold rolled and cold rolled annealed samples are given in Figure 5.24.

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In LD_CR44, the ferrite phase consisted of grains with high angle of misorientation along δ/γ interphase boundaries as well as within the δ-ferrite matrix. Mostly, distorted lattices were developed in the BCC regions during cold rolling. Some portions of BCC phase with low angle of misorientation were witnessed within the austenite islands. These were believed to be the in-grain strain induced martensitic structure. As annealing took place, the BCC phase readily got strain free by favorable recovery process. Sample CR_HT8500.25 more or less fully consisted of strain free BCC grains (blue) just after 15 minutes of annealing. The peak of relative frequency of occurrence increased at very low angle with the increase in annealing time. The longtime annealed sample CR_HT8503 was full of strain free BCC grains and the peak of relative frequency increased to 0.57 at 0.2° misorientation.

Fig. 5.24: The local misorientation maps of BCC phases in cold rolled and cold rolled + annealed samples: (a) LD_CR44, (b) CR_HT8500.25 (850 °C, 15 minutes) and (c) CR_HT8503 (850 °C, 3 hours). The color code bar indicates the grain of particular angle of misorientation (in degrees). 113

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On the other hand, austenite phase manifested a gradual decrease in peak angle of misorientation. The deformation substructure of the austenite phase remained nearly unchanged until the beginning of recrystallization as discovered by Reick et al. (1998). The local misorientation maps of austenite phase in hot rolled and hot rolled annealed samples are provided in Figure 5.25.

Fig. 5.25: The local misorientation maps of austenite phases in cold rolled and cold rolled + annealed samples: (a) LD_CR44, (b) CR_HT8500.25 (850 °C, 15 minutes) and (c) CR_HT8503 (850 °C, 3 hours). The color code bar indicates the grain of particular angle of misorientation (in degrees). The local misorientation maps of austenite phases also revealed the same phenomena while annealing the cold rolled sample. The cold deformed sample LD_CR44 consisted of more number of deformed grains and the most frequent austenite grains were of misorientation angle of 2.3° (green). The deformed grains started to get strain free as the isothermal annealing began to take place. Eventually, the peak relative frequency of misorientation angle shifted to lower value. After prolonged annealing treatment (CR_HT8503), the newly formed tiny and recrystallized γ-grains 114

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were seen throughout the microstructure and they appeared strain free, whereas, some residual stress still remained inside the large grains. This might be the result of presence of internal stress within the lattice as it decreased rapidly up to a certain period of time and then the residual stress remained even after a long exposure at that temperature [Belyakov et al. (2002)]. This finding supported the presence of strained regions within austenite phase.

5.5 Summary & Conclusion The effects of processing parameters on the development of microstructure in DSSs have been studied extensively on 22Cr5Ni3Mo standard duplex grade (SD) as well as on lean duplex steel (LD). The main conclusions that can be drawn from this work are: 1. High temperature annealing promoted ferrite formation. As a result, chromium and molybdenum content decreased and nickel content increased in δ-ferrite phase with increase in annealing temperature. On the other hand, chromium and molybdenum concentration in γ-austenite increased with increase in annealing temperature. 2. Annealing at high temperature followed by rapid quenching resulted in supersaturation of N in δ-ferrite phase of DSS. As a result, excess N segregated at δ/δ grain boundaries and triple points. EDX analysis confirmed that this precipitate was Cr2N. 3. The duplex grades (e.g. SD) enriched with alloying elements, especially with Cr and Mo, were more prone to intermetallic precipitation than the grade (LD) leaner with these elements, owing to faster precipitation kinetics. Upon holding at 900 °C, the as-received standard DSS (SD) resulted in σ-phase formation after only 30 minutes, whereas, the LD material was free from intermetallics even after 5 hours. This finding was in good agreement with the thermodynamic and kinetic calculation. 4. Long-term ageing in the σ-precipitation range promoted 𝛾 → 𝜎 transformation, since, ferrite was almost consumed after 5 hours of ageing treatment. In all the cases, σ-precipitates were enriched with Cr, Mo and depleted with Ni. 5. Ageing lean DSS grade samples at 850-900 °C temperature range encouraged 𝛿 → 𝛾 phase transformation, as expected from the calculated phase fraction diagram.

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However, equilibrium volume % of phases (1:1) could not be achieved even after 5 hours ageing, i.e., phase transformation practically stopped before 5 hours. The bulk hardness of the samples decreased with increasing annealing time. The sample aged at 850 °C for 3 hours consisted of more equiaxed austenite grains (average aspect ratio, AR = 1.56 ± 0.21) than the sample aged at 900 °C for 3 hours (average aspect ratio, AR = 2.13 ± 0.23). The austenite islands transformed from Widmanstätten type to more of equiaxed type as annealing progressed. 6. Hot rolling produced elongated grains along rolling direction and promoted 𝛾 → 𝛿 transformation, however, the amount of transformed γ was independent of degree of the thickness reduction. There was a significant increase in microhardness of both the phases with deformation. Annealing encouraged 𝛿 → 𝛾 transformation and slight decrease in AR value (2.94 → 2.5) of γ-austenite. 7. Hot rolling produced LAGBs which started to move in order to relieve stress during annealing. In case of δ-ferrite, the misorientation angle gradually increased, i.e., LAGBs turned into HAGBs with annealing time. However, the misorientation angle of FCC grain boundaries remained unchanged for a while and sudden characteristic of annealing twin appeared at 60° in the sample that was annealed at 850 °C for 3 h after hot rolling. BCC phase underwent continuous DRX, i.e., recovery and recrystallization, whereas, FCC phase recrystallized through conventional DRX, i.e., nucleation and growth mechanism. BCC phase readily got strain free as annealing started to progress, whereas, FCC phase showed a gradual decrease in peak angle of local average misorientation. Some residual stress remained within the large γ grains even after long-term annealing. 8. Cold rolling produced strain induced lath type αʹ-martensite within γ-austenite, since, C concentration was very low and ε-martensite was metastable, compared with αʹ-martensite. Around 8-9% austenite transformed to martensite after cold deformation. Annealing promoted reverse transformation. Both the bulk hardness and AR value of γ-austenite substantially decreased (430𝐻𝑣 → 250𝐻𝑣 and 8.33 → 1.96 respectively) in cold rolled sample after 3 hours of annealing at 850 °C.

9. Both phases in the cold rolled sample were bound mostly by LAGBs. However, characteristic peak of deformation twins has been observed at 60° in case of FCC. Ferrite possessed higher propensity towards recovery, owing to continuity of the 116

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process which delayed the onset of recrystallization. As a result, even short-time annealing produced HAGBs. The HAGBs formed faster in case of FCC also, due to higher accumulation of strain energy during cold deformation and eventually more inclination towards nucleation and growth, i.e., conventional DRX. Local average misorientation was higher in BCC than in FCC in cold rolled sample, but, lattices were more distorted than hot rolled sample. Some low angles of local average misorientation were also spotted within FCC, which were in-grain strain induced martensitic portions. Annealing promoted quick release of strain energy in BCC, because of its faster overall recrystallization kinetics, mainly to release the stress. In case of FCC, the aspect ratio of the grains started to decrease rapidly as annealing started to progress. Newly formed γ-grains were strain free, whereas, residual strain still remained within large γ-grains even after long-term annealing.

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Chapter 6

EFFECT OF PROCESSING PARAMETERS ON MICROTEXTURE

6.1 Introduction 6.2 Hot rolled and hot rolled-annealed steels 6.3 Cold rolled and cold rolled-annealed steels 6.4 Summary & Conclusion

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6.1 Introduction ‘Texture’ represents the preferred orientation of crystals relative to a fixed reference. The x-ray texture, i.e., macrotexture reveals what volume fraction of the specimen (obtained from the intensity of diffraction by particular planes) has a particular orientation. However, it does not enlighten us how these grains are distributed throughout the material. In the present work, the investigated steel has mixed phases (δ + γ) in its microstructure. The macrotexture would provide gross information about the combined texture of both BCC and FCC phases. Study of microtexture is therefore, more relevant as it provides information about the texture evolution of individual phases. The importance of texture development during processing of metals and alloys on final mechanical properties has been understood in recent years [Herrera et al. (2008), Hamada & Ono (2010)]. Material properties, i.e., strength, ductility, toughness etc. depend considerably on microtexture. To understand the evolution of texture in lean DSS subject to different thermo-mechanical treatments, EBSD study has been performed.

6.2 Hot rolled and hot rolled-annealed steels To determine the texture development during hot rolling and hot rolling-annealing treatment, three samples were taken from LD_HR50 steel and two of those were allowed to anneal at 850 °C for different periods of time (Table 3.3). Inverse Pole Figure (IPF) IPFs were generated by Electron Back Scattered Diffraction (EBSD) technique. The analyses were carried out using HKL CHANNEL FIVE software and the inverse pole figure (IPF) maps of hot rolled and hot rolled-annealed samples are shown in Figure 6.1 for ferrite (BCC) and in Figure 6.2 for austenite (FCC). In case of ferrite (BCC) phase, hot rolling resulted in strong γ-fiber ( || ND) and most of the grains were aligned along the rolling direction (represented by blue), i.e., deviation from this orientation was witnessed in small number of BCC grains. A small fraction of α-fiber ( || RD) was also evidenced and could be identified in annealed samples also. These two are the principal deformation texture within BCC [Ray et al. (1994)]. As the hot rolled samples were subjected to annealing, the misorientation angle between two adjacent δ-grains was also increased and HAGBs were formed readily within 15 minutes of holding at 850 °C. That means continuous 120

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recrystallization was dominated by recovery. Ferrite has high stacking fault energy (SFE) and hence it is more prone to recovery. The orientations of the BCC grains were quite diverse and intensity of γ-fiber got decreased after annealing. Prolonged annealing treatment also promoted cube/rotated cube texture (represented by red). This cube component mainly forms under small strains and shows proneness towards recovery in spite of recrystallization owing to its small in-grain orientation gradients [Ul-Haq et al. (1994)]. Rotated cube component was believed to form upon deformation and recrystallization of cube governed by solidification texture.

Fig. 6.1: Inverse pole figure (IPF) color maps of δ-ferrite phase in hot rolled and hot rolled + annealed samples: (a) LD_HR50, (b) HR_HT8500.25 (850 °C, 15 minutes) and (c) HR_HT8503 (850 °C, 3 h).

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Fig. 6.2: Inverse pole figure (IPF) color maps of γ-austenite phase in hot rolled and hot rolled + annealed samples: (a) LD_HR50, (b) HR_HT8500.25 (850 °C, 15 minutes) and (c) HR_HT8503 (850 °C, 3 h). On the other hand, FCC phase revealed α-fiber ( || ND) in hot rolled sample LD_HR50 (represented by green). This fiber existed even after long term annealing at 850 °C. Cube component was also witnessed in the sample HR_HT8503. To elucidate the texture components more accurately, orientation distribution functions (ODFs) were generated. ODF section φ2 = 45° and φ2 = 0° to 90° were considered in case of BCC and FCC, respectively as it is the most conventional way to represent the texture components.

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Orientation Distribution Function (ODF) ODFs provide information regarding the frequency of occurrence of particular orientations in three dimensional (Euler) orientation space. ODFs of BCC and FCC phases are shown in Figure 6.3 and Figure 6.4, respectively.

Fig. 6.3: φ2=45° sections (Bunge notation) of ODFs of ferrite (BCC) phase in hot rolled and hot rolled + annealed samples: (a) LD_HR50, (b) HR_HT8500.25 (850 °C, 15 minutes) and (c) HR_HT8503 (850 °C, 3 h), (d) standard φ2=45° section for BCC steel showing all texture components [Sk. Md. et al. (2016)].

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Hot rolled sample LD_HR50 revealed both α-fiber and γ-fiber textures. Principal texture of hot rolled BCC phase consisted of {111} component on γfiber, {112} component on α-fiber and {111} component on point where both the fiber lines cut each other. This strongest component was 17.3 times stronger than random texture in BCC grains of the sample. Some weak components on γ-fiber and adjacent region have also been noticed. The texture of BCC phase shifted to rotated cube component {001} from γ-fiber as annealing progressed. This finding supported earlier published result [Badji et al. (2011)]. This rotated cube component was 4.07 times stronger than any random texture. Prolonged annealing promoted cube {100} and rotated cube components. The strongest components within cube fiber were situated in between cube and rotated cube components and the strength was 8.13 times than any random texture. Textures developed during deformation can influence the texture formed after subsequent annealing treatment. Both α-fiber and γ-fiber appear upon deformation of ferrite. Gamma-fiber dominates over α-fiber upon recrystallization of deformed ferrite grains. However, this recrystallization is conventional recrystallization, which involves nucleation and growth of strain-free (recrystallized) grain. In other type of recrystallization which involves pronounced recovery of deformed ferrite grains, sub-grain boundaries gradually turn into high-angle grain boundaries and that leads to recrystallization. This type of recrystallization results in retention of deformed ferrite texture (dominated by α-fiber). It has been reported that melt spinning could result in strong ferritic fiber (α-fiber) which gets weakened after annealing; single {100} component formed [Herrera et al. (2008)]. Another group of researchers showed that the rolling texture component {001} found in an AISI 444 ferritic SS rotated to directions close to {110} after annealing [De Abreu et al. (2006)]. Moreover, in DSSs, the textures of both phases could also be associated by an orientation relation after phase transformation which can take place during annealing. The comparison between ODF intensities of γ-fiber and α-fiber within BCC phase are shown graphically in Figure 6.4a and Figure 6.4b, respectively. Figure 6.4 clearly shows that overall γ-fiber intensity has been decreased and αfiber intensity has been increased as annealing progressed. Sample HR_HT8503 showed that all components within γ-fiber were weak compared to the rolled sample LD_HR50. 124

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On the other hand, α-fiber, specifically, component {112} as well as rotated Cube component {001} and rotated Goss component {110}

got

intensified with annealing treatment. Component {113} appeared after shortterm annealing, but, vanished after long-term thermal exposure at same temperature.

Fig. 6.4: ODF intensities of texture components within (a) γ-fiber and (b) α-fiber in BCC phases of hot rolled sample LD_HR50 and hot rolled-annealed samples HR_HT8500.25 and HR_HT8503. Report suggested that Brass, Copper, S and Goss components were generally observed in deformation texture of austenite, whereas, Cube, Goss and Brass were present in recrystallization texture in FCC phases [Badji et al. (2011)]. The ODF 125

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sections (φ2 = 0° to 90°) of FCC phases are shown in Figure 6.5. The hot rolled sample mainly consisted of MS-Brass {011} texture component (3.8×R) and strong αfiber ( || ND) which is typically a well developed FCC shear texture as reported in the literature [Tóth et al. (1988), Li et al. (2005)]. Weak components were also witnessed in different φ2 sections, e.g. D-Taylor {4 4 11} and Copper {112} texture components at φ2 = 45°, BR-Brass recrystallization {236} component at φ2 = 35° and S-component {123} at φ2 = 65°. Short-term thermal annealing fostered more intense BR-Brass recrystallization {236} component (10.8×R) and, a mixture of D-Taylor {4 4 11} and Copper {112} components (9.49×R) at φ2 = 35° and 45°, respectively, as observed in sample HR_HT8500.25. The α-fiber texture component shifted to Cube component {001}, which is a typical recrystallization texture in FCC phase [Badji et al. (2011)]. Main component after rolling, i.e., MS-Brass {011} texture component was retained after annealing and the intensity increased to 5.42×R. Long term annealing developed more random texture. All texture components developed in the beginning of the recrystallization, remained even after 3 hours annealing, however, their intensities got decreased. Besides, a new Goss component {011} formed in sample HR_HT8503, which is also believed to form during recrystallization of FCC [Badji et al. (2011)]. Several authors suggested that the exact proportion of deformation and recrystallization texture components is controlled by the stacking fault energy of FCC phase [Smallman & Green (1964)]. Other group of researchers reported that the same depended on stored energy, i.e., strain accumulated during deformation [Chowdhury et al. (2005)]. Extensive study on the MS-Brass texture in hot rolled DSSs was carried out by several authors [Hirsch et al. (1988), Ray (1995), Vercammen et al. (2004)] who usually delineated it through various possible mechanisms owing essentially to the low stacking fault energy of the material, such as deformation twinning and shear banding. An additional Goss component was found in DSSs after hot rolling, which is often attributed to the elevated deformation temperature which tends to generate extra shear strain, predominantly near the surface layer [Kang et al. (2007)]. A partial retention of deformation texture components during annealing was usually observed in 2 cases if – (i) stored strain energy is low, and (ii) microstructure changes during annealing is

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dominated by recovery (mostly in ferrite). However, recrystallization texture is more readily expected in FCC as this phase undergoes recrystallization instead of recovery.

Fig. 6.5: φ2=0° to 90° sections (Bunge notation) of ODFs of austenite phase in hot rolled and hot rolled + annealed samples: (a) LD_HR50, (b) HR_HT8500.25 (850 °C, 15 minutes) and (c) HR_HT8503 (850 °C, 3 h), (d) standard φ2=0° to 90° sections for FCC steel showing all texture components. 127

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Interfaces with Kurdjumov-Sachs (K-S) orientation relationship The Kurdjumov-Sachs (K-S) relationship is one of the most frequently cited orientations associated with δ-γ interphase boundaries other than NishiyamaWassermann (N-W) relationship [Kurdjumov & Sachs (1930)]. The ledge mode of misfit accommodation is extended to {111}FCC || {l10}BCC interfaces with the Kurdjumov-Sachs (K-S) orientation relationship. Unlike N-W relationship, closepacked directions are also parallel for K-S, i.e., FCC || BCC [Van der Merwe & Shiflet (1994)]. The maps of hot rolled and hot rolled-long term annealed samples highlighting phase boundaries with K-S orientation relationship are shown in Figure 6.6.

Fig. 6.6: K-S orientation maps of hot rolled sample (a) LD_HR50 and hot rolled-long term annealed sample (b) HR_HT8503 (850 °C, 3 h), color code represents angle of deviation from K-S. The boundaries with K-S orientation are complex semi-coherent interfaces. The dislocations in the {111}FCC || {l10}BCC interfaces are observed to be sessile [Knowles & Smith (1982)]. As δ and γ co-deform, softer δ phase deforms predominantly and that develops high stress concentration at the δ/γ interface. Interphase boundary sliding can release that stress concentration, but, sliding is difficult in K-S boundaries due to its semi-coherent nature. Therefore, interphase boundary cracking occurs if it undergoes further deformation. The austenite phase can precipitate near K-S orientation relationship with at least one side of the ferrite grain. Normally, this side does not grow into the ferrite grain because of low energy grain boundary. Nevertheless, the side

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without a fixed crystallographic orientation relationship grows into the ferrite grain [Herrera et al. (2008)]. Interphase boundaries deviate from ideal K-S orientation relationship upon deformation. It could easily be understood that annealing for 3 hours resulted in extensive number of newly developed δ/γ interphase close to K-S orientation which were not seen in hot rolled sample. It can be inferred that long term annealing is detrimental for further deformation, since, cracks can generate from these interphase during forming or bending operation. Twinning It has been reported that microstructural observations could substantiate the absence of recrystallization twins within austenite grains [Iza-Mendia et al. (1998)]. Twinning has been observed in recrystallized austenite in DSSs [Iza-Mendia et al. (1996)]. A typical FCC twin system is patently a superposition of two sets of diffraction patterns that are symmetrical to each other with respect to {111} planes [Zhang et al. (2003)]. Figure 6.7 shows how the amount of twins increased after annealing the hot rolled sample.

Fig. 6.7: Maps of hot rolled sample (a) LD_HR50 and hot rolled-long term annealed sample (b) HR_HT8503, highlighting twin boundaries (red) within austenite grains (yellow). Annealing for 3 hours resulted into several annealing twins which were not witnessed in sample LD_HR50. These twins are high energy and high angle boundaries, and show characteristic peak at 60° angle of misorientation as shown in Fig. 5.16b.

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6.3 Cold rolled and cold rolled-annealed steels To elucidate the texture development during cold rolling and cold rolling-annealing, three samples were taken from LD_CR44 steel specimen and two of those were annealed at 850 °C for different periods of time (section 3.3.3 and 3.3.4) to study the chronological texture evolution in cold rolled and cold rolled-annealed steels. Inverse Pole Figure (IPF) Like hot rolled and hot rolled-annealed steels, EBSD technique has been employed to generate IPF maps of cold rolled and cold rolled-annealed steels and analyses were performed in HKL CHANNEL FIVE. The IPF color maps of δ (BCC) and γ (FCC) phases of steel samples LD_CR44, CR_HT8500.25 and CR_HT8503 are presented in Figure 6.8 and Figure 6.9, respectively.

Fig. 6.8: Inverse pole figure (IPF) color maps of δ-ferrite phase in cold rolled and cold rolled + annealed samples: (a) LD_CR44, (b) CR_HT8500.25 (850 °C, 15 minutes) and (c) CR_HT8503 (850 °C, 3 h).

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IPF map of cold rolled sample LD_CR44 revealed a mixture of strong γ-fiber ( || ND) (blue) and cube/rotated cube (red) texture component. The γ-fiber of hot rolled sample LD_HR50 was retained even after cold rolling and shifted from α-fiber. This is in agreement with the previous published results [Ray et al. (1994)]. However, some authors have suggested that a typical BCC cold rolling texture was α-fiber [Hamada & Ono (2010)]. To know the origin of cube/rotated cube texture component, ODF maps needed to be generated. As annealing started to progress, Goss texture (green) developed within BCC at the expense of γ-fiber. Long term annealing produced more random texture with no particular component strengthening, although, retainment of γ-fiber has been witnessed in sample CR_HT8503.

Fig. 6.9: Inverse pole figure color maps of FCC phase in cold rolled sample (a) LD_CR44 and cold rolled-annealed samples (b) CR_HT8500.25 and (c) CR_HT8503, annealed at 850 °C for 15 minutes and 3 hours, respectively. In case of FCC phase, cold rolled sample LD_CR44 consisted of mostly alphafiber texture component ( || ND) (green). Annealing resulted in weakening of alpha-fiber and strengthening of component on gamma-fiber || ND (blue). The randomness of texture increased with annealing time. 131

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Orientation Distribution Function (ODF) The ODF maps of BCC and FCC phases of cold rolled and cold rolled-annealed samples are presented in Figure 6.10 and Figure 6.11, respectively. ODF at φ 2=45° section of BCC phase in sample LD_CR44 revealed strong γ-fiber with main components {111} (~7.5×R), {111} (~5.5×R) and a mixture of ̅̅̅̅3> and {111} (~4.5×R). Similar observation was reported by Herrera {111} and {332} were also et al. (2008). The transition components {554} component shifts towards {554} and Report suggests that {332} during subsequent rolling. At large strains, there is a net rotation then to {111} to the {111} position [Ray et al. (1994)]. Since, applied from the {111} and {111} (~7.5×R), {111} (~5.5×R) and a mixture of {111} (~4.5×R). Brass component {110} (in the austenite phase of {111} and {111} (~7.5×R), {111} (~5.5×R) and a mixture of {111} (~4.5×R). Brass component {110} (in the austenite phase of {111}