Effect of Microstructure and Crystallographic Texture

 

Effect of Microstructure and Crystallographic Texture on Mechanical Properties of Modified 9Cr-1Mo Steel    

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

  of

 

Doctor of Philosophy  

by

 

Arya Chatterjee (Roll No: 12MT92P02)    

Under the supervision of

 

Dr. Debalay Chakrabarti and

Prof. Rahul Mitra  

 

 

Department of Metallurgical and Materials Engineering Indian Institute of Technology Kharagpur May, 2018  

© 2018 Arya Chatterjee. All rights reserved.

 

             

Dedicated to

My beloved parents & brother  

 

 

 

APPROVAL OF THE VIVA-VOCE BOARD  

   

--/--/2018  

   

             

Certified that the thesis entitled EFFECT OF MICROSTRUCTURE AND CRYSTALLOGRAPHIC TEXTURE ON MECHANCIAL PROPERTIES OF MODIFIED 9Cr-1Mo STEEL submitted by ARYA CHATTERJEE to the Indian Institute of Technology, Kharagpur, for the award of the degree Doctor of Philosophy 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 DSC)

(Member of the DSC)

(Member of the DSC)

 

               

(Supervisor)

(Supervisor)

 

               

(External Examiner)

(Chairman)

 

                               

CERTIFICATE  

   

This is to certify that the thesis entitled Effect of Microstructure and Crystallographic Texture o n Mechanical Properties of Modified 9Cr-1Mo Steel, submitted by Arya Chatterjee 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 Doctor of Philosophy of the Institute.

           

            _________________________ 

(Dr. Debalay Chakrabarti) Associate Professor Dept. of Metallurgical and Materials Engineering, Indian Institute of Technology Kharagpur-721302, India    

Date:

 

 

     ________________________ 

(Prof. Rahul Mitra) Professor Dept. of Metallurgical and Materials Engineering, Indian Institute of Technology Kharagpur-721302, 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.

       

 

(Arya Chatterjee)

 

ACKNOWLEDGEMENTS I take this opportunity to express my deep sense of gratitude to Dr. Debalay Chakrabarti and Professor Rahul Mitra for their invaluable guidance and constant inspiration throughout the PhD program. I have benefited by learning a lot from them both academically and personally. Their incomparable enthusiasm and sense of commitment has inspired me. In addition, I really feel that this work could not be completed without their friendly treatment towards me during my stay at Indian Institute of Technology Kharagpur. I sincerely convey my thanks and appreciation to the members of the Doctoral Scrutiny Committee, Professor K. K. Ray Professor G. G. Roy, Dr. Sumatra Mondal and Professor Santanu Dhara (School of Medical Science and Technology) and all other professors for their comments and suggestions whenever required to enrich my PhD work. I am highly grateful to Board of Research in Nuclear Sciences, India for financially support during early years of my PhD work. I like to express deep sense of gratitude to Dr. Arun K. Bhaduri and Dr. Aniruddha Moitra of Indira Gandhi Centre for Atomic Research, Kalpakkam for their help and support at each and every stage of this project. I would like to express my special thanks to Prof. Laszlo S. Tóth, LEM3 Laboratory, University of Lorraine and Dr. Cyril Cayron, Laboratory of Thermomechanical Metallurgy-PX Group, EPFL for the insightful discussions with them related to the some of my PhD work. I thank Mr. Dilip Chakraborty, Mr. Pradip Sarkar, Mr. Tinku Thomas, Mr. Biswarup Kar, Mr. Prasanta Das, and all the staff members of Dept. of MME, Mr. Niloy Bhoumick, Mr. Tapas Pal, Mr. Ranadhir Basu and all the staff members of Central Research Facility, IIT Kharagpur for their various types of assistance during experimental works. I am thankful to all my friends and research colleagues from IIT Kharagpur, especially discussions with Dr. Abhijit Ghosh and Dr. Kaustav Barat enriched my knowledge. I consider myself to be extremely fortunate to have friends like Kaku, Khana, Abhi, Baka, Chulke, Johny, Sourav Da, Anupam, Arka, Pavan, Hasan, Basir, Hura, Bibhu, Gunti, Pan, Chhoru during my stay in IIT Kharagpur. It has been a great pleasure for me to spend time and have fun with them. Last but not the least, I deeply acknowledge my indebtedness to my parents, Dada, Jagrity Di, Partha, Soumita and other well-wishers and friends for their good wishes and encouragement.

____________________ (Arya Chatterjee)  

vi  

   

List of Symbols  

Symbol  *f

 fd

Description Cleavage fracture stress Dynamic flow stress

 *f ( 85)

Cleavage fracture stress at -85°C temperature.

 *f (Tgy )

Cleavage fracture stress at or below Tgy temperature.

 yd

Dynamic yield stress.

g

Grain level stress Aggregate level stress Peak stress in the plastic zone Dislocation density Plastic strain-rates at individual grain Plastic strain-rates at aggregate level Kronecker delta, which is one for valid data and zero for censored data.



ˆ 

 g 

i 

0 γp

Shear strain rate Reference slip rate Energy required for plastic work prior to the creation of new fracture surface

γmm

Effective surface energies for matrix-matrix interface

γpm γs µ

Effective surface energies for particle-matrix interface Energy required to create a unit area of new fracture surface shear modulus Reference critical resolved shear stress

ca

1, , 2 Eulers angle in Bunge notation representing the orientation of any crystal in sample reference frame. Cleavage angle between adjacent crystals  CL

CLmeas

Cleavage angle between adjacent crystals measured on crack plane

CLPr oj

CCL

Calculated projection angle of  CL on measuring plane Twist angle Initial crack length of Charpy impact specimen Increment in crack length Orientation of austenite in regard to ith variant of martensite Burger’s vector Thickness of Charpy impact specimen Initial ligament width of Charpy impact specimen Thickness of the tested Charpy impact specimens. Thickness of the predicted Charpy impact specimens. cleavage plane normal ( type) of crystal

CF

Plastic constraint factor

θtwist a0 Δa Ai b B b0 B1 B2

vii  

List of Symbols CC CS CV

Plane-normal in crystal reference frame Plane-normal in sample reference frame Charpy impact energy of the V-notch samples

d D Davg Deff d pl

Displacement value determined from instrumented impact test Grain size Average grain size Effective grain size Plastic displacement

DPacket  90

Size corresponds to 90th percentile of carbide size distribution

DCarbide‐90

Size corresponds to 90th percentile of carbide size distribution Young’s modulus Plastic energy absorbed by the specimen up to the maximum load

E

Emp gexp  G

h

ab

J J0 J0.2 J0.2t JId KIC KJd Kmin Kpcf K Jd (1) K Jd (1 inch ) K Jd (2)

M

MVi

Orientation of martensite, experimentally obtained Unit vector of grain boundary plane normal Latent hardening matrix J-integral Dynamic elastic-plastic fracture toughness for Δa = 0 mm Dynamic elastic-plastic fracture toughness for Δa = 0.2 mm Dynamic elastic-plastic fracture toughness for Δa = 0.2 mm from the intersection of blunting line with the J-R curve Dynamic elastic-plastic fracture toughness Plane strain static fracture toughness Elastic plastic fracture toughness obtained from JId A constant fracture toughness value= 20 MPa-m1/2 Plastic constraint factor KJd of first (10 mm thick) sample. Equivalent of KJd for 1-inch (25.4 mm) thick specimen. KJd of second (25.4 mm thick) sample. Grain interaction tensor orientation matrix of ith number of crystallographic variant of martensite

Pgy

Dynamic strain hardening exponent Load value obtained from instrumented impact test Crack arrest load Maximum load estimated from load-displacement curve of instrumented impact test below Tgy temperature Dynamic flow load, estimated from load-displacement curve of instrumented impact test General yield load estimated from load-displacement curve

Pi Pmax t T T0 T28J

Brittle crack initiation load Maximum load Time value obtained from instrumented impact test Charpy impact test temperature Reference temperature for fracture transition Transition temperature corresponds to Charpy impact energy absorption of

n P Pa Pf

Pfd

viii  

List of Symbols T4kN TDBTT Tgy

T0dyMSP TVi TNR R2 W Y

Temperature corresponds to the arrest-load in Charpy impact test as 4 kN Transition temperature corresponds to ductile-brittle transition temperature Temperature indicating brittleness transition temperature, where Pgy = Pmax Dynamic reference temperature following modified Schindler’s procedure Transformation matrix of ith martensitic variant Recrystallization stop temperature Root mean square Width of Charpy impact specimen precipitation hardening effects of alloying elements

     

ix  

 

x  

List of Abbreviations Abbreviation Ae1 Ae3 Ar1 Ar3 BCC Carbide-90 CVN DBTT DRX EBSD ECD EDS EGS EPFM FATT GOS HAB HEM HR ITT J-R LAB LEFM LSE m-m MPS MSP ND NT Packet-90 PAG PAG-90 PAGS PCF p-m RD S.D SEM TD TEM T-L TMCR USE USE UTS YS

Description Equilibrium lower critical temperature Equilibrium upper critical temperature Lower critical temperature during cooling Upper critical temperature during cooling Body centered cubic 90th percentile size of the coarse carbide Charpy V-notch Ductile-brittle transition temperature Dynamic recrystallization Electron back scattered diffraction Equivalent circular diameter Energy dispersive spectroscopy Effective grain size Elastic–Plastic Fracture Mechanics Fracture appearance transition temperature Grain orientation spread High-angle boundaries Homogeneous Effective Medium Hot-rolled Impact transition temperature J-integral and Resistance to fracture curve low-angle boundaries Linear Elastic Fracture Mechanics Lower shelf energy obtained from Charpy impact transition curve Matrix-matrix Martensitic packet size Modified Schindler procedure Normal direction Normalized and tempered 90th percentile size of martensitic packets Prior-austenite grain 90th percentile size of prior-austenite grains Prior-austenite grain size Plastic constraint factor Particle-matrix Rolling direction Standard deviation Scanning electron microscope Transverse direction Transmission electron microscope Transverse-longitudinal thermo-mechanical controlled rolling Upper-shelf energy Upper shelf energy obtained from Charpy impact transition curve Ultimate tensile strength Yield strength

 

xi  

                                           

xii  

List of Tables Table 2.1: Energy Scenario in India. Table 2.2: Threshold criterion used for the determination of effective grain size for different microstructures in steel. Table 2.3: Different Plastic Constraint Factor (PCF) values reported in the literature. Table 2.4: Identification of the particular mechanism of creep through knowledge of the creep parameters, n, p and QC. Where Qgb is the grain boundary diffusion activation energy and Q is the lattice diffusion activation energy Table 3.1: Chemical composition of the investigated samples, (wt.%). Table 3.2: Sample codes with processing details along with the section numbers where these are mentioned. Table 5.1: Microstructural parameters, fraction of low-angle boundaries and the effective grain sizes (estimated from EBSD analysis) and cleavage facet sizes (measured from fractographic study) of the As-received and hot-rolled samples. Table 5.2: Tensile properties of the investigated samples. Table 5.3: Upper shelf energy (USE) and impact transition temperatures obtained from the instrumented Charpy impact testing of the investigated samples. Table 6.1: Microstructural parameters of the as-received and reheated samples. Fraction of low-angle boundaries and the effective grain sizes were measured from the EBSD analysis. Cleavage facet sizes were measured from the fractographic study. Table 6.2: Tensile properties obtained from the tension test on the investigated samples. Table 6.3: Upper shelf energy (USE) and impact transition temperatures obtained from the instrumented Charpy impact testing of the investigated samples. Table 7.1: Different transition temperatures, upper shelf energies, and cleavage fracture stresses at Tgy and at -85°C of the investigated samples. Table 7.2: Different microstructural parameters used to correlate with different impact transition properties of the investigated samples. Table 7.3: Adjusted R2 values of the different microstructural parameters used to * correlate with cleavage fracture stress,  f ( 85) of the investigated samples.

Table 8.1: Voce hardening parameters for simulating flow curves of shear stress-shear strain of different samples at strain rate of 10 s-1. Table 8.2: Different martensitic grains, effective grain sizes for threshold misorientation of 15° considering angle-axis pair (EGS) and cleavage angles (EGS{100}), and dynamic impact properties (USE and DBTT) of the investigated samples. xiii  

List of Tables Table 8.3: Few examples of interacting martensitic boundaries and cleavage crack given in cracked microstructures and cleavage facet. Table 9.1: The creep properties of the normalized samples. dy Table A1: T0 MSP related data of the investigated steel samples.

Table A2: Kurdjumov-Sachs (K-S) orientation relationships of 24 crystallographic variants (V), and their corresponding transformation matrix (T) that correlates them with their parent austenite. Table A3: Minimum angles between adjacent slip planes for different KS inter-variants boundaries (IVB) considering only {110} slip planes (A{110}), and considering both {110} and {112} slip planes (A(both)).

 

xiv  

List of Figures Fig. 1.1:

(a) Optical micrograph showing martensitic microstructure of 9Cr-1Mo steel. Prior-austenite grain boundary (PAGB) and martensitic packet boundaries are delineated with red and yellow colour lines, respectively; (b) Transmission electron micrograph indicates parallel martensitic laths within a typical martensitic packet [13]. Arrangement of different units of martensitic microstructure at different length scale has been schematically represented in (c) [14].

Fig. 1.2:

Schematic representation of different martensitic boundaries and presence of different kinds of carbides (M23C6 and MX types) along those boundaries.

Fig. 2.1:

Application of 9Cr-1Mo steel in fuel tubes for nuclear reactor application.

Fig. 2.2:

Charpy impact curves for Sandvik HT9 steel in unirradiated condition and after irradiation to 10 and 17 dpa at 365ºC [3].

Fig. 2.3:

Schematic showing the typical rolling schedule and the associated microstructural changes for thermomechanical controlled rolling (TMCR) of steels [50].

Fig. 2.4:

Deformation and transformation textures develop in ferritic-martensitic steel [52].

Fig. 2.5:

Different texture components of ferritic-martensitic steel at φ2=45° section of the Euler’s space.

Fig. 2.6:

Stereographic projection of {002} poles for each of the 24 K-S variants in (111)γ single crystal matrix [63]. The orientation relationships of the symbols and code numbers are shown at right side in tabular form.

Fig. 2.7:

(a) Experimental (002) pole figure for transformation texture of martensite obtained from heavily rolled austenite in Fe-25·7Ni alloy [63], and (b-d) simulated (002) pole figures of transformation texture of martensite using the BP, AS, and TS models, respectively [61].

Fig. 2.8:

Fracture surface study of martensitic grade steel showing (a) cleavage facet forms during brittle transgranular fracture (nucleated from Cr23C6 particle, as indicated by arrow), and (b) ductile fracture, indicated by the presence of micro-voids nucleated from spheroid carbide particles (shown by arrows).

Fig. 2.9:

Schematic diagram showing variation in yield stress and fracture stress with temperature. TGY denotes the general yield temperature.

Fig. 2.10: Schematic diagram showing different steps involved in cleavage fracture: Step 1: nucleation of sharp microcrack nucleates at some microstructural feature, Step 2: propagation of microcrack across particle-matrix interface, Step 3: propagation of microcrack across matrix-matrix boundaries [67]. Fig. 2.11: Characteristic points on the Charpy impact transition curves. Solid line represents impact energy transition curve and dotted line represents fracture appearance transition curve. xv  

List of Figures Fig. 2.12: Characteristic points on schematic load-time profile in the transition region. Fig. 2.13: Schematic diagram showing the different fracture mechanisms operating at different temperatures determined considering the variation in yield strength (σ0), the peak stress in the plastic zone ((σˆ) and the cleavage fracture stresses considering particles (σpm) and grains (σmm) [70]. Fig. 2.14: Variation in cleavage fracture stress (σF, MPa) with D-1/2 showing the increase in σF with the decrease in grain size, D (mm) [71]. Fig. 2.15: Influence of grain size (mm) and carbide thickness (µm) on 27 J-ITT [73]. Fig. 2.16: (a) A cleavage facet from TMCR steel on which orientation imaging (OIM) was carried out; (b) Corresponding OIM image of one cleavage facet from TMCR steel (four different grains are marked); (c) Inverse pole figure of the orientation of crystals detected on fracture surface [17]. Fig. 2.17: Schematic diagram showing a crystal with an arbitrary orientation with respect to the fracture plane of the sample. RD, TD and ND are the rolling direction, transverse direction and normal direction, respectively.[86]. Fig. 2.18: A typical creep curve indicates three different regions: the primary, secondary and the tertiary creep region [49]. Fig. 2.19: Illustration of the effect of stress and temperature on creep behavior of any Material. Fig. 2.20: Coarsening of lath width after long term creep exposure in a ferriticmartensitic steel. (a): Virgin material, and (b) material in crept condition [113]. Fig. 3.1:

Schematic diagrams showing (a) single pass deformation schedules, (b) multi pass deformation schedules.

Fig. 3.2:

(a) Schematic representation of 15 pass deformation schedule applied to the ‘As-received’ steel using Gleeble 3500® thermo-mechanical simulator, and (b) the corresponding mean flow stress (MFS) of each pass plotted against inverse of deformation temperature to determine the point of change in slope, which represents recrystallization stop temperature (TNR).

Fig. 3.3:

Schematic diagram showing the heat-treatment and rolling schedules used in the present study.

Fig. 3.4:

Schematic diagram showing the heat-treatment schedule used is the present study.

Fig. 3.5:

Schematic representation of different thermo-mechanical processing routes with different colors being used for samples deformed at different hot-rolling (HR) temperatures.

Fig. 3.6:

Typical dimension and orientation of samples for microstructural study, EBSD study, texture study, tensile testing and Charpy impact testing with xvi

 

List of Figures respect to the original rolled plate Fig. 3.7:

Schematic diagram showing the dimensions of sub-size tensile Specimen.

Fig. 3.8:

Schematic diagram showing the dimensions of tensile creep Specimen.

Fig. 3.9:

Schematic diagram showing the dimensions of full size Charpy impact sample.

Fig. 3.10: Schematic diagram showing the deviation in estimation of DBTT ( ± ΔDBTT) and USE (± ΔUS) values from tanh curve fitting. Fig. 4.1:

(a) Flow curves for the single pass deformed samples and (b) the variation in mean flow stress with strain for the multi-pass deformed samples.

Fig. 4.2:

Optical micrographs of single pass (SP) deformed samples. Sample codes are mentioned on the images.

Fig. 4.3:

Optical micrographs of multi pass (MP) deformed samples. Sample codes are mentioned on the images.

Fig. 4.4:

EBSD micrographs showing high angle boundaries (black coloured) and low angle boundaries (red coloured) in single pass deformed samples. Sample codes are given on the images.

Fig. 4.5:

EBSD micrographs showing high angle boundaries (black coloured) and low angle boundaries (red coloured) in multi pass deformed samples. Sample codes are given on the images.

Fig. 4.6:

Variation in (a) average grain size and (b) grain aspect ratio as the function of finish deformation temperature in single pass and multi pass deformed samples. Abbreviations: PAG: Prior austenite grain size, Eff: Effective grain size, SP: Single pass, MP: Multi pass.

Fig. 4.7:

Typical EBSD inverse pole figure maps with respect to normal direction (ND-IPF) of selected samples (sample codes mentioned) used for the determination of ODF. Colour legends for IPF maps are inserted.

Fig. 4.8:

A typical Gleeble-deformed sample. RD, TD and ND stand for deformation direction (i.e. simulated rolling direction), transverse direction and normal direction, respectively.

Fig. 4.9:

(a-d) Orientation distribution function (ODF) at φ2=45° sections (Bunge notation) of single pass (SP) deformed samples. Samples codes are mentioned on the maps; (e) φ2=45° ODF sections showing the ideal orientations as observed in ferritic steels (BCC). Abbreviations: Rot-C: Rotated cube; Rot-Goss: Rotated Goss.

Fig. 4.10: Orientation distribution function (ODF) at φ2=45° sections (Bunge notation) of multi pass (MP) deformed samples. Samples codes are mentioned on the maps. Fig. 4.11: Transmission electron micrographs showing precipitates (indicated by xvii  

List of Figures arrows) and their corresponding EDS and SAED analysis in the deformed (a) SP-1273, (b) SP-1173 (c) MP3-0.8 and (d) MP3-0.6 samples of modified 9Cr-1Mo steel. Fig. 4.12: Volume fraction of different precipitates predicted using Thermo-Calc® thermodynamic software as the function of temperature (centre) and the compositional variation within the respective precipitates (mentioned on the images) with temperature. Fig. 4.13: The recrystallization-precipitation-time-temperature (RPTT) diagram for (a and b) Nb(C, N) for applied strain of 0.2 and 0.8; (c) RPTT diagram for VN for 0.2 and 0.8 applied strain. RPTT diagrams are calculated following Dutta and Sellars [178] and Medina et al. [179,180]. Abbreviations: Ps: Precipitation start time, Pf: Precipitation finish time, Rs: Recrystallization start time and Rf: Recrystallization finish time. Fig. 5.1:

(a) Optical and (b) SEM micrographs of the As-received steel.

Fig. 5.2:

Optical micrographs of the hot-rolled samples: (a) 1050-HR, (b) 1000-HR, (c) 950-HR and (d) 875-HR.

Fig. 5.3:

SEM micrographs of hot-rolled samples: (a, b) 1050-HR, (c, d) 1000-HR, (e, f) 950-HR and (g, h) 875-HR; (a, c, e, g) before tempering and (b, d, f, h) after tempering treatment. Ferrite regions are indicated by arrows.

Fig. 5.4:

Transmission electron micrographs of (a) As-received, (b) 1050-HR, (c) 1000-HR, (d) 950-HR and (e) 875-HR samples, (f) dark field image of the precipitates and (g) corresponding SADP analysis identifying the precipitate as Nb(C,N). Precipitates are marked by arrows.

Fig. 5.5:

The EBSD image quality maps showing the different misorientation boundaries in: (a) As-received, (b) 1050-HR, (c) 1000-HR, (d) 950-HR and (e) 875-HR samples.

Fig. 5.6:

Cumulative distributions of the boundary misorientation angles measured (using EBSD) across the different boundaries present in the investigated samples.

Fig. 5.7:

Orientation distributions presented in the 2=45 sections of the Euler space as obtained from the macro-texture study of (a) As-received, (b) 1050-HR, (c) 1000-HR, (d) 950-HR and (e) 875-HR samples.

Fig. 5.8:

Representation of (a) RD-fiber, (b) TD-fiber and (c) ND-fiber components in the investigated steel samples. RD, ND and TD represent rolling direction, transverse direction and normal direction, respectively.

Fig. 5.9: Engineering stress-strain curves obtained from the tension tests of the investigated samples. Fig. 5.10: (a) Volume fraction of precipitates formation predicted using Thermo-Calc® thermodynamic software as the function of temperature, (b) enlarged view of xviii  

List of Figures micro-alloying precipitate formation. Fig. 5.11: Precise lattice parameters of As-received and hot-rolled specimens considering bcc structure: (a) before the tempering treatment and (b) after the final tempering treatment. Fig. 5.12: Charpy impact transition curves obtained from the Charpy impact testing of the investigated samples from 9Cr-1Mo steel. Fig. 5.13: Correlation between (a) volume fraction of cleavage planes parallel to main fracture plane (i.e. {001}//RD-ND plane) and DBTT of the investigated samples and (b) volume fraction of slip planes of the crystals parallel to the shear plane of the samples (i.e. {110}+{112})//shear plane) and USE of the investigated samples. Fig. 5.14: SEM fractographs of (a, f) As-received, (b, g) 1050-HR, (c, h) 1000-HR, (d, i) 950-HR and (e, j) 875-HR samples tested at the (a-e) upper shelf regime (test temperature of +80C) and (f-j) lower shelf regime (test temperature of -100C). Some of small voids are indicated by black arrows, while large voids are indicated by red arrows. Fig. 6.1:

Optical micrographs of As-received and reheated samples of 9Cr-1Mo steels: (a) As-received, (b) 950-Reheat, (c) 1025-Reheat and (d) 1100-Reheat.

Fig. 6.2:

SEM micrographs of (a, b) As-received steel and (c-h) reheated samples. Micrographs of the reheated samples are taken (c-e) before tempering and (fh) after tempering treatment.

Fig. 6.3:

The EBSD image quality maps showing the different misorientation boundaries in the investigated samples: (a) As-received sample, (b) 950Reheat sample, (c) 1025-Reheat sample and (d) 1100-Reheat sample.

Fig. 6.4:

Cumulative distributions of the boundary misorientation angles measured (using EBSD) across the different boundaries present in the investigated samples.

Fig. 6.5:

The φ2=45º sections of the Orientation Distribution Function (ODF) as obtained from the macro-texture study on all the investigated samples: (a) As-received sample, (b) 950-Reheat sample, (c) 1025-Reheat sample and (d) 1100-Reheat sample.

Fig. 6.6:

Transmission electron micrographs of (a) As-received sample, (b) 950Reheat sample, (c) 1025-Reheat sample and (d) 1100-Reheat sample; (e) dark field image the precipitates and (f) the SADP analysis identifying the precipitate circled in (e) as (Nb,V)(C,N). Precipitates are indicated by arrows in (a-e).

Fig. 6.7:

Vickers macro-hardness of the investigated samples measured at 20 kg load. The error bar represents the standard deviation in measured hardness values with respect to the average hardness values. Hardness of reheated samples xix

 

List of Figures was measured before and after the reheating treatment. Fig. 6.8:

Volume fractions of precipitates in the investigated steel as predicted from Thermo-Calc® software at different reheating temperatures. (a) Solid line represents the equilibrium volume fraction of Cr23C6 precipitates and dotted line represents the volume fraction of Cr23C6 precipitates expected under actual (non-equilibrium) reheating condition. (b) Equilibrium volume fractions of Nb(C,N) and V(C,N) precipitates as predicted from ThermoCalc®.

Fig. 6.9:

Lattice parameters of As-received and reheated specimens as measured from the XRD analysis: (a) lattice parameters measured before the final tempering treatment and (b) lattice parameters measured after the tempering treatment.

Fig. 6.10: Engineering stress-strain curves obtained from the tension tests of the investigated samples from 9Cr-1Mo steel. Fig. 6.11: Charpy impact transition curves obtained from the Charpy impact testing of the investigated samples from 9Cr-1Mo steel. Fig. 6.12: SEM fractographs of (a, e) As-received sample, (b, f) 950-Reheat sample, (c, g) 1025-Reheat sample and (d, h) 1100-Reheat sample tested at the (a-d) upper shelf regime (test temperature of +40C) and (e-h) lower shelf regime (test temperature of -130C); (i) the energy dispersive spectroscopy (EDS) analysis of the particle indicated by arrow in (e). Fig. 7.1:

Typical load-displacement plot of the as-received sample tested at 0°C showing general yield load (Pgy) and maximum load (Pmax).

Fig. 7.2:

Typical variation in load values (Pmax and Pgy) obtained from loaddisplacement plots of 1000-HR sample as the function of test temperature. Brittleness transition temperature (Tgy) is the temperature corresponding to the intersection point of two 2nd order polynomial fitted curves representing the variation in Pgy with temperature and Pmax with temperature.

Fig. 7.3:

Variation of different impact transition temperatures as the function of (a) reheating temperatures for normalized samples and (b) finish rolling temperatures for hot-rolled samples.

Fig. 7.4:

Correlation between T28J and T0MSP fro the investigated samples. Continuous

dy

red-line indicates existing established relation for low strength ferritic steel and dotted lines represent standard deviation. Fig. 7.5:

(a) Correlation coefficient, adjusted R2 of all the microstructural parameters in absolute and logarithmic values while correlating with different transition temperatures, and (b) 90th percentile size of martensite packet size (Packet90) in logarithmic values showed best fit with the transition temperatures of investigated samples.

Fig. 7.6:

Relation between 90th percentile carbide size (Carbide-90) and cleavage xx

 

List of Figures fracture stress estimated at -85°C (  *f ( 85) ) for the investigated samples. Fig. 7.7:

Correlation between upper shelf energy (USE) and (a) 90th percentile carbide size (Carbide-90) and (b) effective grain size of the investigated samples.

Fig. 7.8:

Cross-sectional view of the Charpy impact specimens tested at -196°C illustrating secondary cleavage cracks showing: (a, b) crack deflection or retardation at prior austenite grain boundaries and martensite packet boundaries (marked with ‘A’ and ‘B’, respectively). Packet boundaries unable to alter crack propagation path is indicated by ‘C’, whereas, crack stops within a packet (‘D’), (c) Cleavage crack initiation from carbide cracking is indicated by arrows.

Fig. 8.1:

(a) Schematic representation of K-S variants surrounded by prior-austenite grain boundary (PAGB). Abbreviation, SBB: sub-block boundary; BB: block boundary; PB: packet boundary. (b) All possible misorientation (angle-axis pair) and cleavage angles between different types of boundaries and their distributions. (c) Minimum misorientation angles considering angle-axis pair and cleavage angles, and cleavage twist angles considering the grain boundary plane normal (GBPN) as specimen axis for every combination of KS variants. Legends of (c): Red, Green and Blue colours indicate combination of variants that forms sub-block, block and packet boundaries, respectively, and the table below shows the location of different types of angles in tabular form of (c).

Fig. 8.2:

Typical optical micrographs of the investigated samples for (a) 875HR950NT, and (b) 1050HR-1100NT samples.

Fig. 8.3:

Textures of the microstructures (IPF maps) for investigated samples. (a) 875HR-950NT, (b) 875HR-1025NT, (c) 875HR-1100NT, (d) 950HR950NT, (e) 950HR-1025NT, (f) 950HR-1100NT, (g) 1000HR-950NT, (h) 1000HR-1025NT, (i) 1000HR-1100NT (j) 1050HR-950NT, (k) 1050HR1025NT, (l) 1050HR-1100NT. 1: grain boundaries (GB) considering a minimum 15° misorientation angle based on angle-axis pair, and 2: GB considering minimum cleavage angle as 15°. (m) IPF legend and micrometer scale bar.

Fig. 8.4:

Charpy impact transition curves of the investigated samples indicating absorbed impact energy as function of temperatures.

Fig. 8.5:

Typical fractographs of 950HR-1100NT specimens broken at (a) brittle (110°C), (b) ductile (+45°C) and (c) mean transition temperature (-30°C) regimes showing cleavage facets, dimples (voids) and quasi-cleavage features, respectively. Facets and voids are indicated by red and yellow arrows, respectively.

Fig. 8.6:

Correlation of effective grain size considering angle between {100} cleavage planes (EGS{100}) with ductile-to-brittle-transition temperatures (DBTT) of xxi

 

List of Figures the samples. Fig. 8.7:

Typical EBSD inverse pole figure (IPF) maps showing secondary cleavage crack propagation in martensitic structure across different boundaries between martensitic variants for different investigated samples.

Fig. 8.8:

Typical examples of secondary cleavage crack propagation and their deviation at different martensitic boundaries between martensitic variants.

Fig. 8.9:

EBSD study of cleavage facets in fracture surface.

Fig. 8.10: Schematic representation of cleavage crack propagation across three different martensitic variants. Fig. 8.11: Distribution of angles between slip planes for all combination of martensitic variants considering (a) {110} as active slip system, and (b) both {110} and {112} as active slip systems. Fig. 8.12: Optical micrographs of adiabatic shear band (ASB) zones in 1000HR1100NT samples fractures at high temperature (+75°C). Rotated and straight ASBs are indicated by indicated by red and yellow arrows, respectively, whilst cracks at rotated ASBs are shown by black arrows. Fig. 8.13: Optical micrographs showing prior-austenite grain (PAG) structures of (a) 1000HR-950NT (S1), (b) 875HR-1100NT (S2), (c) 1050HR-1100NT (L1) and (d) 1000HR-1100NT (L2) samples. (e) PAG size distributions of these samples. Fig. 8.14: IPF maps showing adiabatic shear bands (ASBs) adjacent to ductile shear cracks of (a) S1, (b) S2, (c) L1 and (d) L2 samples (a-d counter clockwise sequence). IPF legend is given in inset. Fig. 8.15: 1: Initial, 2: experimentally obtained, and 3: simulated textures (i.e. pole figures of (110), (111) and (112), and ODF at φ2=45°) of (a) S1, (b) S2, (c) L1 and (d) L2 samples. Fig. 8.16: Variation in average grain rotation within adiabatic shear band regions of the investigated samples. The grain rotation is defined as the rotation angle around the rotation axis of the initial ellipsoids from the initial into the final position. Fig. 8.17: Typical Euler angle map of 875HR-1100NT sample with reconstructed prioraustenite grain structure shown in inset. The identified variants in one of the prior-austenite grain, ‘A’ (orientation: φ1=90.3°, φ=104.1° and φ2=180.8°) are indicated in the microstructure. Fig. 8.18: Distributions of inter-variants boundaries for different investigated samples. (a) 875HR-950NT, (b) 875HR-1025NT, (c) 875HR-1100NT, (d) 950HR950NT, (e) 950HR-1025NT, (f) 950HR-1100NT, (g) 1000HR-950NT, (h) 1000HR-1025NT, (i) 1000HR-1100NT (j) 1050HR-950NT, (k) 1050HR1025NT, (l) 1050HR-1100NT. xxii  

List of Figures Fig. 8.19: Similar secondary cleavage crack regions as that shown in Fig. 8.7. (a) Elastic modulus maps show variation in modulus (in terms of colours) at the locations of significant crack deviations, and (b) strain distributions adjacent to crack path obtained from Kernel Average Misorientation (KAM) maps. Colour legends are shown in upper-left corners. Fig. 8.20: Number of active slip systems per grain within ASB regions of different investigated samples along the strain path. Fig. 8.21: Fraction of slip activity for {110} and {112} slip systems along the strain path within the ASB regions of (a) S1, (b) S2, (c) L1 and (d) L2 samples. Fig. 8.22: Schematic representation of a single crystal subjected to loading. The crystal initially deforms primarily with slip primarily in slip system, S1. Lattice rotations (α) occurs such that slip system S2 rotates toward the loading axis which eventually leads to the conjugate system to be active in S2` state. Therefore, slip activity in conjugate slip systems promotes lattice rotation. Fig. 8.23: Intra-grain orientation deviations within adiabatic shear band regions are shown by Grain orientation spread (GOS) maps of (a) S1, (b) S2, (c) L1 and (d) L2 samples. Fig. 8.24: Typical transmission micrographs from ASB regions of (a) L2 and (b) L1 showing fragmented grains in bright field (BF) and dark field (DF) mode with selected area diffraction (SAD) pattern in inset, and of (c) L2 and (d) S1 indicating rarely visible fragmented grains, sub-cells and high dislocation density within laths shown by yellow, blue and red arrows, respectively. (e) Examples of dislocations crossing GBs in S1 sample, higher magnification BF image of A and B zones from (d). Fig. 8.25: (a) Typical HREM image of micro-voids in the formation of crack near GB triple points in ASB region of L2. Enlarged view shows disclination dipoles and partial disclination indicated by red and yellow ovals, respectively, (b) presence of similar disclination dipoles in L1, but without formation of crack or void. Fig. 8.26: Typical transmission electron micrographs showing of presence of M23C6 in (a) HAADF-STEM image and (b) dark-field image, and MX type precipitates can be seen in (b) dark-field image as well as (c) highmagnification bright field image. M23C6 and MX precipitates are indicated by red and yellow arrows, respectively. Fig. 8.A1: Shear stress-shear strain curves used for determination of Voce hardening parameters. Line represent experimental data, while points are simulated ones. Fig. 9.1:

Schematic diagram showing normalized sample subjected to three different creep test temperatures.

Fig. 9.2:

(a) Creep strain vs. time to rupture plots of all the investigated samples. Creep xxiii

 

List of Figures curves of normalized samples tested at 550 °C (b), 600 °C (c) and 650 °C (d). Minimum creep rate and duration of creep tests of the investigated samples are shown in (e) and (f), respectively. Fig. 9.3:

Grain orientation spread (GOS) maps of (a) 950NT-550, (b) 1025NT-550, (c) 1100NT-550, (d) 950NT-600, (e) 1025NT-600, (f) 1100NT-600, (g) 950NT-650, (h) 1025NT-650 and (i) 1100NT-650 samples, where GOS < 1 (blue in color) indicates the recrystallized grains. .

Fig. 9.4:

(a) Recrystallization grain area fractions and (b) ratio of primary creep strain to total creep strains of the investigated normalized samples as function of creep test temperatures.

Fig. 9.5:

Sub-surface creep damage evolution near ruptured surfaces of maps of (a) 950NT-550, (b) 1025NT-550, (c) 1100NT-550, (d) 950NT-600, (e) 1025NT600, (f) 1100NT-600, (g) 950NT-650, (h) 1025NT-650 and (i) 1100NT-650 samples. Slip markings in (b) and (d), grain boundary slidings in (d-f), and creep cavities in (g-i) are indicated by arrows.

Fig. 9.6:

Transmission electron micrograph of (a) 950NT-550 indicates that sub-grain boundary (shown by dotted yellow line) is not retarded by Cr23C6 precipitates and crossed the frontal part of the precipitates. (b) In 1025NT-550, Cr23C6 type precipitates are anchoring the sub structural martensitic lath boundaries, whilst (c) in 1100NT-550 some of the Cr23C6 precipitates are unable to hold the martensitic lath boundaries during creep deformation. (b) and (c) are STEM-HAADF images. Interactions between MX types of precipitates and dislocations indicates Orowan bowing in (d) 1025NT-600 and (e) 1100NT600 specimens. (f) Incoherent Cr23C6 precipitate-dislocation interaction at the detachment end of the particle showing attractive force between dislocation and particles in 1025NT-650 sample. (d-f) all are weak beam dark filed (WBDF) images with g = [011] having zone axis [100], one representative diffraction pattern is shown in (g).

Fig. 9.7:

Orientation distribution function (ODF) maps at the φ2 = 45° section of Euler space for (a) 950NT-550, (b) 1025NT-550, (c) 1100NT-550, (d) 950NT-600, (e) 1025NT-600, (f) 1100NT-600, (g) 950NT-650, (h) 1025NT-650 and (i) 1100NT-650 samples. (j) φ2=45° ODF section showing the ideal orientations as observed in ferritic steels (BCC).

Fig. 9.8:

Inverse pole figure (IPF) maps along the rolling direction (RD) for (a) 950NT, (b) 1025NT, (c) 1100NT, (d) 950NT-550, (e) 1025NT-550, (f) 1100NT-550, (g) 950NT-600, (h) 1025NT-600, (i) 1100NT-600, (j) 950NT650, (k) 1025NT-650 and (l) 1100NT-650 samples.

Fig. 9.9:

Distribution of relative fraction of boundary misorientation angles of the investigation samples (a) before creep test, and after creep tests at (b) 550 °C, (c) 600 °C and (d) 650 °C.

Fig. 9.10: Image quality (IQ) maps obtained from the EBSD scans of (a) 950NT-550, xxiv  

List of Figures (b) 1025NT-550, (c) 1100NT-550, (d) 950NT-600, (e) 1025NT-600, (f) 1100NT-600, (g) 950NT-650, (h) 1025NT-650 and (i) 1100NT-650 samples. Fig. 10.1: Ductile-to-brittle transition temperature (DBTT), upper-shelf energy (USE) and time to creep rupture of all the differently processed modified 9Cr-1Mo steel samples investigated in the present study.

                                                                   

xxv  

 

xxvi  

 

Abstract A modified 9Cr-1Mo steel received in normalized and tempered condition has been separately subjected to different normalization treatments varying the austenitization temperatures (1100C / 1025C / 950C) and different hot-rolling treatments varying the finish rolling temperatures (1050C / 1000C / 950C / 875C). A combination of both these processing was also carried out by hot-rolling at different temperatures and normalizing each hot-rolled plate from different temperatures. The amount of deformation applied on each hot-rolled plate was constant having true strain of ~ 0.7. The processed samples were finally tempered at 750C and tested for tensile and Charpy impact properties following standard procedures. A substantial improvement in upper-shelf energy (USE) by 20 J and reduction in ductile-to-brittle transition temperature (DBTT) by 20C has been noticed in normalized samples, in comparison to the As-received plate, especially for the sample normalized from 1025C. Although hot-rolling increased the strength significantly (by 100 MPa), it reduced the ductility by more than 10% and increased the DBTT by more than 25C. Among the hot-rolled samples, the plates rolled at 1050C showed the best combination of high USE (comparable to the As-received steel) and low DBTT. The results have been analysed considering the effect of normalization and rolling temperatures on the parameters related to the microstructure and crystallographic texture of the steel, such as, prior-austenite grain size and martensitic packet size, ferrite fraction, effective grain size, fraction of the low-angle boundaries, fraction of {100} cleavage planes parallel to the main fracture plane and the fraction of {110 and 112} slip planes along the 45 to the main fracture plane (i.e. along the maximum shear stress plane). ‘Effective grain size’ represents the crystallographic unit over which cleavage crack propagates in an uninterrupted fashion. For the study of macro- and micro-texture, X-ray diffraction based texture goniometer and electron back-scattered diffraction (EBSD) technique were used, respectively. The Impact properties of the investigated samples were represented in terms of USE, DBTT, 28J impact energy transition temperature, general yield temperature (Tgy), dy

blunt-notched dynamic fracture toughness (Kjd) and reference temperature ( T0MSP ). Kjd was determined following the modified Schindler’s procedure for the same strain rate as that of the standard Charpy impact tests. Correlation between different transition temperatures and their conservativeness were analysed in view of the fracture safe design. Finally, microstructural parameters were related to the impact transition temperatures based on the microstructural and fractographic studies. The effect of hierarchical martensitic microstructure having different structural units of varying length scales (i.e. lath, sub-block, block, packet and prior-austenite grain) on the micro-mechanism of deformation and fracture have been elucidated by studying the propagation of cleavage cracks and the formation of shear cracks within the investigated samples using EBSD technique and visco-plastic self-consistent (VPSC) polycrystalline plasticity model. The study indicates strong influence of certain crystallographic variants on the c xxvii  

 

leavage crack propagation. The ‘martensitic block’ was found to be the ‘effective grain’ controlling the impact toughness at low temperatures, where cleavage fracture dominates. Dynamic fracture at high temperatures, on the other hand, was found to be dictated by cracking along the shear bands, evolution of which depend on the size and distribution of the prior-austenite grains. An attempt has also been made to study the effect and evolution of microstructure and texture of the normalized samples during creep deformation at three different temperatures (550C / 600C / 650C). The sample normalized at 1025C show better creep resistance at the low temperature (550°C) and the high temperature (650°C) creep regimes. The creep properties are correlated with the microstructural parameters and their stability, precipitate-dislocation interaction and crystallographic texture. Keywords: 9Cr-1Mo Steel, Martensite, Crystallographic texture, Thermomechancial processing, Charpy impact testing, Tensile testing, Creep testing, Effective grain, Adiabatic shear band.        

xxviii  

Contents Subject Title Page Certificate of Approval Certificate Declaration Acknowledgements List of Symbols List of Abbreviations List of Tables List of Figures Abstract Contents Chapter 1 Introduction 1.1 General background 1.3 Objectives 1.4 Organization of the thesis Chapter 2 Literature review 2.1 Requirement of steel for high temperature application 2.2 Variants of 9Cr-1Mo steel and applications 2.3 Processing and properties 2.3.1 Thermomechanical Processing 2.3.1.1 Development of crystallographic texture during processing of ferritic-martensitic steel: 2.3.1.2 Variant selection of martensite during thermomechanical processing: 2.3.2 Normalizing-tempering treatment 2.4 Fracture Mechanism 2.4.1 Ductile fracture 2.4.2 Brittle fracture 2.5 Instrumented Charpy impact testing 2.6 Effect of microstructure on fracture behavior along the impact transition regime 2.7 Effect of crystallographic texture on impact transition behavior 2.8 A review on studies on impact transition behaviour of 9-12 wt% Cr-Mo steel 2.9 Creep Mechanism 2.9.1 Standard Creep equations 2.9.1.1 Effect of stress and temperature 2.9.1.2 Effect of microstructure 2.9.2 Identifying the mechanisms of creep 2.10 A short review on creep behaviour studies of 9-12%Cr-Mo steel 2.11 Summary Chapter 3 Experimental details 3.1 Materials 3.2 Processing and heat treatment schedules 3.2.1 Study on the effect of hot deformation schedule on the microstructure 3.2.2 Study and effect of hot-rolling to vary the microstructure and properties xxix  

Page no i iii iv v vi vii xi xiii xv xxvii xxix 1-7 1 4 5 9-39 9 10 12 12 13 14 17 18 18 18 20 23 26 30 33 35 35 36 37 38 39 41-50 41 41 41 43

3.3.3 Study on the effect of normalization temperature on the 44 microstructure and properties 3.2.4. Study on the combined effect of hot-rolling and normalization 44 temperature on the microstructure and properties 3.3 Microstructural characterization 46 3.4 Texture study 48 3.5 Hardness testing 48 3.6 Uniaxial tensile testing and tensile creep testing 48 3.7 Instrumented Charpy impact testing 49 3.8 Fractography 50 Chapter 4 Effect of microalloy precipitates on microstructure and 51-68 texture of hot-deformed modified 9Cr-1Mo steel 4.1 Introduction and objective 51 4.2 Results 52 4.2.1 Study of hot flow behavior 52 4.2.2 Microstructural study 52 4.2.3. Micro-texture study on hot-deformed samples 56 4.2.4. Precipitation study in hot-deformed samples 59 4.3 Discussion 62 4.3.1. Effect of precipitates on microstructure and texture 62 4.4 Summary 67 Chapter 5 Effect of hot-rolling temperature on the ductile-brittle 69-85 transition behaviour modified 9Cr-1Mo steel 5.1 Introduction and objective 69 5.2 Results and Discussion 70 5.2.1 Characterization of microstructure 70 5.2.2. Characterization of inclusions and precipitates 72 5.2.3. Characterization of boundary misorientation 74 distribution and texture 5.2.4 Tensile properties of the investigated samples 76 5.2.5 Charpy impact properties of the investigated samples 78 5.2.6 Effect of grain size, packet size and ferrite fraction 79 on impact properties 5.2.7 Effect of precipitate particles on impact properties 80 5.2.8 Effect of crystallographic texture on impact properties 81 5.2.9 Fractographic studies on the impact tested specimens 82 5.3 Summary 84 Chapter 6 Effect of normalization temperatures on ductile-brittle 87-104 transition behaviour of modified 9Cr-1Mo steel 6.1 Introduction and objective 87 6.2. Results and Discussion 88 6.2.1 Characterization of microstructure 88 6.2.2 Characterization of boundary misorientation 90 distribution and texture 6.2.3 Characterization of inclusions and precipitates 92 6.2.4 Hardness and tensile testing of investigated samples 93 6.2.5 Charpy impact testing of investigated samples 97 6.2.6 Effect of grain size and packet size on impact 98 transition behavior 6.2.7 Effect of precipitate particles on impact toughness 99 xxx  

6.2.8 Effect of macro-texture on impact toughness 101 6.3 Summary 103 Chapter 7 Dynamic fracture behaviour of thermo-mechanically 105-122 processed modified 9Cr–1Mo steel 7.1 Introduction and objective 105 7.2 Calculation procedures 7.2.1 Determination of Tgy temperature from load-time 106 data processing: 7.2.2 Estimation of JId from CVN specimens 107 7.3 Results and Discussion 108 dy 7.3.1 Determination of dynamic reference temperatures ( T0MSP ) 108 7.3.2 Comparative study between transition temperatures 110 7.3.3 Effect of microstructural parameters on impact 113 transition temperatures (ITT) 7.3.4 Effect of microstructural parameters on cleavage 115 fracture stress 7.3.5 Effect of microstructural parameters on upper shelf energy 117 7.3.6 Cleavage fracture micro mechanism 119 7.4 Summary 120 Chapter 8 Combinations of hot-rolling and subsequent 123-164 normalization treatment on impact fracture behavior and the influenced of martensitic hierarchical microstructures on micro-mechanism of deformation and fracture 8.1 Introduction and objective 125 8.1.1 Background 8.1.2 Dynamic fracture of martensitic steel at low temperature 125 8.1.3 Plastic deformation and fracture of martensitic steel 126 in upper shelf region. 8.2 Modelling approaches 129 8.2.1 Identification of crystallographic variants in martensite 129 8.2.2 Modelling dynamic fracture associated with 133 deformation plasticity 8.3. Results 135 8.3.1 Materials before deformation 135 8.3.2 Charpy impact tests 137 8.3.3 Martensitic variants and low temperature dynamic fracture 139 8.3.4 Hierarchical martensitic microstructures and high 144 temperature dynamic deformation 8.3.5 Evolution of crystallographic texture within 146 adiabatic shear bands 8.3.6 Rotational effects within adiabatic shear bands 150 8.4. Discussions 151 8.4.1 Effect of crystallographic variants on transition 151 temperatures 8.4.2 Elasticity and micro-plasticity of cleavage crack propagation 155 8.4.3 Rotational phenomenon within the adiabatic shear bands 156 8.4.4 Microstructural damage within adiabatic shear bands 159 8.5. Summary 164 xxxi  

Chapter 9

Effect of normalization treatment on creep strength at different creep temperatures 9.1 Introduction and objective 9.2. Results and Discussion 9.2.1 Creep behaviour of investigated samples 9.2.2. Effect and evolution of microstructure in relation to creep 9.2.3 Effect and evolution of precipitates in relation to creep 9.2.4 Evolution of crystallographic texture after creep 9.2.5 Evolution of boundary distribution after creep 9.3 Summary Chapter 10 Conclusions and future scope of work 10.1 Conclusions 10.2 Future scope of work Appendix Bibliography

   

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165-182 165 166 166 168 171 174 177 180 183-188 183 188 189-194 195-217

Chapter 1 Introduction

   

   

Chapter 1 1.1

General background The selection of 9Cr-1Mo steel for Fast Breeder Reactor (FBR) applications is

primarily based on a good combination of mechanical properties, high thermal conductivity, low thermal expansion coefficient, and good resistance to stress-corrosion cracking in water–steam systems compared to austenitic stainless steels [1-31]. The initial FBR programs used austenitic stainless steel as wrapper material in the core subassembly, but due to high irradiation induced void swelling and irradiation creep in austenitic stainless steel, which caused dimensional instability and core distortion, ferritic-martensitic grade 9Cr-1Mo steel has been adopted for these applications in FBR [1]. In the austenitic stainless steels, void swelling produces an increase in length and diameter of the wrappers and causes the subassemblies subjected to the gradients of fast neutron flux and temperature, to bow and to interact with its neighbours or core restraint structure. This alloy also exhibits good weldability and microstructural stability over very long periods of exposure to high-temperature service conditions. The major areas of concern/interests/investigation on 9Cr-1Mo steel for last two decades are the following [1-31]:

 Type IV cracking of weld [7, 16, 19-20, 23-25],  Creep rupture [6-9],  Thermal fatigue [7-12],  Creep-fatigue interaction [7, 9, 12],  Oxidation resistance [7, 17], and  Stress corrosion cracking [7, 14] The fracture resistance, especially in the ductile-to-brittle transition temperature (DBTT) regime is a matter of concern for reactor pressure vessel steels. Irradiation during the long service condition severely affects the dynamic fracture resistance of steel, by increasing the DBTT, [1, 3, 7, 15, 29-31]. As a result even at room temperature the steel can become brittle and prone to cracking under impact loading either inside the reactor or during external handling of fuel tubes [2–4]. Therefore, the fracture toughness of the steel in unirradiated condition has to be as good as possible, with DBTT as low as possible and upper-shelf energy (USE) has to be as high as possible.

1

Chapter 1  

1.2

Factors affecting impact toughness Martensite in 9Cr-1Mo steel has a hierarchical microstructure comprised of

different structural units of varying length scale like lath, sub-block, block, packet and prior-austenite grains (PAG), Fig. 1.1. As the cleavage crack can nucleate easily from any of the numerous carbide particles (Fig. 1.2), the DBTT of ferritic/martensitic steel primarily depends on the resistance offered by the microstructural barriers, i.e. the microstructural boundaries, to the cleavage crack propagation [5–8] . Now, the boundaries separate the martensitic units are different in nature and have different misorientation angles across them [9–11]. Earlier studies showed the beneficial effect of refining the PAG size and martensitic packet size in lowering the DBTT as their boundaries effectively retard / divert the cleavage crack propagation [5–8]. Low temperature along with high strain rate results in brittle fracture for body-centred cubic (BCC) structure like ferritic/martensitic steel [12].

Fig. 1.1: (a) Optical micrograph showing martensitic microstructure of 9Cr-1Mo steel. PAG and martensitic packet boundaries are delineated with red and yellow colour lines, respectively; (b) Transmission electron micrograph indicates parallel martensitic laths within a typical martensitic packet [13]. Arrangement of different units of martensitic microstructure at different length scale has been schematically represented in (c) [14].

Fig. 1.2: Schematic representation of different martensitic boundaries and presence of different kinds of carbides (M23C6 and MX types) along those boundaries. Extensive amount of research has been carried out on ferritic/martensitic steels subjected to dynamic fracture (especially impact loading) at cryogenic temperatures. 2

Chapter 1 Klueh et al. and Tchizhik et al. [5–8] mentioned the importance of PAG size and martensitic packet size in deciding the toughness of 9Cr-1Mo steel at low temperature regime. Several reports suggested that finer austenite grain size and finer (martensitic) packet size generated by various thermo-mechanical processing resulted in better fracture resistance of bainitic / martensitic steels [15–20]. Localized deformation and fracture characteristics of cleavage crack in martensitic steel depends on the ability of different martensitic boundaries to resist the crack propagation. Previous investigations indicated different opinions in this regards. Hughes et al. [21] indicated cleavage crack propagation across PAG and lath boundaries, whilst others [22–24] reported PAG boundaries as effective in crack retardation. Most of the earlier research showed martensitic packet boundaries as effective barrier to cleavage crack propagation [25–28]. These investigations either correlated the size of cleavage facets in transgranular fracture modes with the packet size and block size [19,21–23] or estimated the ‘effective grain’ size considering only the high-angle boundary (HAB) [17,18,25–27]. In these literatures, various types of boundaries (e.g. packet boundaries and block boundaries) are differentiated in terms of angle of misorientation (angle-axis) exist in those boundaries, which may be misleading for characterizing their effectiveness in cleavage crack retardation as cleavage crack passes through {001} type of planes in neighbouring crystals for BCC materials [24,25,29,30]. Moreover, the role of block boundaries are still unclear from this regard. However, Guo et al. [31] considered the theoretical aspects of Bain variants and their possible effect on cleavage fracture. As per the present knowledge of the authors, there are very few studies available that dealt with direct observation of cleavage crack path in view of crystallography of martensite [27,32]. As per the above discussion, it has become clear that the effect of microstructural parameters especially PAG size, martensitic lath, block and packet sizes on the toughness of modified 9Cr-1Mo steel has not been fully understood from the earlier investigations. Therefore the microstructural features originating from prior processing routes dictate the properties of this steel. 9Cr-1Mo steel is generally used in normalized and tempered condition [9, 11, 20, 27]. Microstructural parameters not only depend on the heat-treatment schedule but also on the prior deformation history [3236]. Till date the effect of hot-deformation on the microstructural parameters and its effect on fracture resistance of 9Cr-1Mo steel is not clear. However, heat treatment and

3

Chapter 1  

hot-deformation plays critical role in changing the grain structure and precipitation of M23C6 (primarily Cr23C6) and MX (mostly Ti(C,N), Nb(C,N) and V(C,N)) type precipitates in modified 9Cr-1Mo steel. However, that aspect has not been checked experimentally in 9Cr-1Mo steels by comparing different microstructures, having different PAG sizes, martensitic packet sizes and / or block sizes. Moreover, the role of these precipitates, present in the modified 9Cr-1Mo steel, on the development of microstructure and texture should be known in order to understand the microstructuretexture-property correlation in this grade of steel. However, proper relation between brittle fracture-creep properties and microstructure for the present steel should be established to improve the ductile-brittle transition behavior for avoiding irradiation embrittlement by tailoring microstructure, texture and precipitation without much affecting its creep properties to be used in power plant application.

1.3

Objectives The major objectives of the present study on the processing-microstructure-

property correlation in modified 9Cr-1Mo steel are listed below: (i)

Developing different microstructure and texture by varying thermomechanical processing schedules in terms of difference in amounts of deformation and inter-pass time, and studying the role of precipitation in 9Cr1Mo steel.

(ii)

Altering martensitic microstructure to develop different morphology and sizes for different martensitic units, such as prior-austenite grain (PAG), martensitic packet, block and lath, and different crystallographic textures by applying hotrolling at different rolling temperatures. The effect of these different microstructures on the strength and impact properties will be investigated.

(iii) Varying the sizes of different microstructural units in martensitic microstructure using different normalization temperatures, and studying their effects on the mechanical properties like tensile properties and impact toughness. (iv) Developing a method for the determination of dynamic fracture toughness, which will be cost effective and more reliable in terms of rapid assessment of toughness during industrial application as compared to the existing methods.

4

Chapter 1 (v)

Studying the combined effect of rolling temperature and normalizing temperature on the impact toughness properties in order to achieve optimum combination of properties. Based on this study, understanding the deformation and fracture micro-mechanisms both at room temperature and very low temperature.

(vi)

Studying the effect of normalizing temperature on the microstructure, texture and precipitation and finally on the creep resistance. Based on this study, understanding the microstructural and textural changes during creep deformation at different temperatures.

1.4

Organization of the thesis

The thesis is divided into several chapters as described below: Chapter 1: Introduction. Chapter 2: Literature review. Chapter 3: Experimental details. In Results and Discussions part all the above objectives are discussed in detail: Chapter 4

(First objective)

Effect of microalloy precipitates on the microstructure and texture of hotdeformed modified 9Cr-1Mo steel. Microalloying elements like Nb and V are added to modified 9Cr-1Mo steel to ensure excellent creep resistance by the formation of fine MX precipitates during tempering treatment. Industrial processing of modified 9Cr-1Mo steel often develops deformed and elongated prior-austenite grain structure, which can be detrimental from property point of view. The effect of Nb and V rich precipitates on the evolution of microstructure (and texture) in hot deformed steel has been studied by performing deformation simulations using Gleeble® 3500. Chapter 5

(Second objective)

Effect of hot-rolling temperatures on the impact toughness and transition behaviour of modified 9Cr-1Mo steel. As-received steel has been subjected to different hot-rolling treatments at different rolling temperatures based on above, at or below the recrystallization stop 5

Chapter 1  

temperature (as known from earlier chapter). The amount of deformation is kept constant for all the hot-deformation treatments. Depending upon the temperatures of deformation, the size and morphology of martensitic grain (PAG, martensitic packet, blocks and lath) structures, the extent of strain-induced precipitations, and the crystallographic textures have been changed. The effect of these different microstructure and texture on the tensile and Charpy impact properties have been assessed in the present chapter. Chapter 6

(Third objective)

Effect of normalization temperatures on the impact toughness and transition behaviour of modified 9Cr-1Mo steel. The temperature domain within which applied deformation would result in good combination of equiaxed austenite grain structure and beneficial texture is established. Moreover, with the knowledge of recrystallization stop temperature for modified 9Cr1Mo steel determined from the previous chapter, the current material has been exposed to different normalization treatments by varying the temperature of normalization. The normalization treatment alters significantly the martensitic grain structure from its asreceived condition as expected. Moreover, the amount of precipitation as well as texture formation is also changed. The role of these different microstructure, texture and the precipitation on the impact toughness and tensile behavior of 9C-1Mo steel have been analyzed in this chapter. Chapter 7

(Fourth objective)

Dynamic fracture behaviour of thermo-mechanically processed modified 9Cr– 1Mo steel. Impact properties obtained earlier from different heat treated, and hot-rolled specimens were represented in terms of USE, DBTT, 28J energy transition temperature, general yield temperature, blunt-notched dynamic fracture toughness (Kjd) and reference temperature. Kjd was determined following the modified Schindler’s procedure with the same strain rate as that of Charpy impact tests instead of considering reduced strain-rate data. Correlation between different transition temperatures and their conservativeness were analyzed in view of fracture safe design. Finally in this chapter, microstructural parameters were related to the impact transition temperatures based on microstructural and fractographic studies. 6

Chapter 1 Chapter 8

(Fifth objective)

Combined effect of rolling temperature and the subsequent normalizing temperatures on impact properties, and the influence of hierarchical martensitic microstructure on the micro-mechanism of deformation and fracture in modified 9Cr-1Mo steel. As-received 9Cr-1Mo steel was hot-rolled at different temperatures applying same amount of deformation, and each hot-rolled specimens were normalized using different austenitizing temperatures to change the martensitic grain structure and texture of the thermo-mechanically processed samples. These samples having tempered martensitic microstructures were impact tested over a range of temperatures. The effect of hierarchical martensitic microstructure with different structural units of varying length scales (i.e. lath, sub-block, block, packet and PAG) on the micro-mechanisms of deformation and fracture at different temperature regimes. The ‘effective grain’ for brittle fracture resistance in low-carbon martensitic steel has been established, and the role of martensitic grains in high temperature is also discussed in great details in this chapter. Chapter 9

(Sixth objective)

Effect of normalization temperature and the consequent microstructural parameters on the creep resistance of the modified 9Cr-1Mo steel. Use of different normalizing temperatures altered the microstructure and precipitation of the modified 9Cr-1Mo steel with respect to the as-received condition (as discussed in earlier chapters). Effect of microstructure and texture on the creep properties of these samples such as, minimum creep rate, creep strain and rupture time have been investigated in this chapter. This investigation has been done to have an intial idea about the heat treatment schedule which is optimum for better impact as well as creep property point of view. Chapter 10: Conclusion and future scope of work.

7

 

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8

 

Chapter 2 Literature review

   

 

   

Chapter 2 2.1

Requirement of steel for high temperature application Indian economy is developing at a faster rate as compared to the previous

decades to be a major economic power in the world in years to come with the growth rate envisaged to be around 10% of the GDP (Gross Domestic Product). With the increase in economic activity, the demand for electricity is expected to rise and by 2020 the electricity demand in this country may touch 300 GWe. Among the total amount of production of electricity, presently coal-fired thermal power plants are the prime source in India, followed by hydroelectric power, contribution of gas, oil and other renewable energy sources and nuclear power plants, Table 2.1 [33]. Table 2.1: Energy Scenario in India. Fuel

Installed Capacity (MWe) *31.09.2017

Coal

1,93,427

Hydro

44,653

Gas

25,185

Renewable Energy Resources

58,303

Nuclear

6,780

Oil

838

Total

329,226

As part of the efforts of preserving the environment it is necessary to reduce of the CO2 emissions from power plants. This can be done by increasing the plant efficiency. Therefore, it was of prime interest for the researcher to be engaged in developing new steels capable of sustaining higher stresses and temperatures required for higher efficiency of power plants. The thermal efficiency of power plant is to a large extent dependent on the operating temperature and pressure of the steam cycle, which are limited by the properties of the materials used in power plant components [34,35]. One of the key materials for power plant application is 9-12% Cr steels. Apart from other desirable properties, improvement of creep strength for this grade of material from operating condition of 180 bar / 530-540 °C to 300 bar / 600 °C reduced the CO2 emission ~ 30% over the last decades [36].

9

Chapter 2  

2.2

Variants of 9Cr-1Mo steel and applications Modified 9Cr-1Mo steel used in the present study comes under the variants of

9-12 wt% Cr steels. There are many potential applications for the 9-12 wt% Cr steels, but the single largest use is in the power generation industry. Specifically, they have found use in super heaters and reheaters tubing, boilers, main steam pipes, bolting and turbine blades and rotors in fossil fuelled power plants [37]. With the application of these steel in power plant components the efficiency of power generation is expected to increase significantly [38]. Therefore, 9Cr grade steel is presently considered as a potential candidate material for futuristic ‘ultra super critical’ power plant in India. The design and production of 9-12 wt% Cr martensitic/ferritic steels first begun in 1912 with 12 wt% Cr steel having 2-5 wt% Mo [39]. Over the next 50 years, this type of steel steels was increasingly used for applications in power plants [35,40]. In mid 1960s, V was added and 12 wt% CrMoV steels were introduced having higher creep strength because of solution hardening and contribution from precipitation of M23C6 carbides [41]. In the late 1970s the P91 steel (modified 9Cr-1Mo) was developed for manufacturing of pipes and vessels for fast breeder reactors [41,42]. The improvement of creep strength for this steel as compared to the earlier 12 wt% CrMoV steels is owing to the formation of thermally stable V and Nb rich precipitates. Moreover, lower Cr content of about 9 wt% formed tempered martensitic microstructure, which also contributes to the higher creep strength [43,44]. For the application of fuel cladding in fast fission reactors the austenitic steels were initially considered as the candidate material till mid-seventies [45]. Due to the high volumetric expansion of the austenitic steels during irradiation (swelling), their applications for cladding and piping as well as for structural components were challenged by the swelling phenomenon [46]. Among other steels and alloys, ferritic and the tempered martensitic steels were also considered and appeared as good alternatives to the austenitic ones during mid-seventies. These steels were then introduced into the international fusion material programs in the early eighties. By now, ferritic martensitic grade P91 steels are one of leading candidate material for fusion structural applications as well, particularly for fuel tubes, Fig. 2.1.

10

Chapter 2

Fig. 2.1: Application of 9Cr-1Mo steel in fuel tubes for nuclear reactor application. The main challenge regarding the applications of the tempered martensitic steels in nuclear power plant application is related to their intrinsic ductile to brittle transition. The fracture mode changes from the high-temperature microvoid coalescence to the low-temperature quasi-cleavage one for this grade of materials [47]. As a consequence of neutron irradiation, the transition temperature between the ductile and brittle regime is shifted towards higher temperatures, limiting the use of these steels in applications to fusion first wall and blanket structures [48]. Within the framework of the fusion material development program, the fracture properties have been mainly investigated with Charpy impact tests. However, Charpy impact testing is actually a very simplistic approach to assess the ductile to brittle transition. The technique consists of determining the so-called ductile-to-brittletransition temperature (DBTT) by measuring the total energy necessary to fracture a V-notched specimen (see for instance [49]). Recently, Charpy impact testers are instrumented so that the load-time curve can be obtained and evaluated. A typical Charpy curve for a tempered martensitic steel is presented in Fig. 2. 1 [3]. The effect of irradiation on the Charpy curve is also depicted in Fig. 2.2. In general, irradiation induces a shift of the DBTT to higher temperature and a decrease of the upper shelf energy.

Unirradiated 

Irradiated 

Fig. 2.2: Charpy impact curves for Sandvik HT9 steel in unirradiated condition and after irradiation to 10 and 17 dpa at 365ºC [3].

11

Chapter 2  

2.3

Processing and properties

2.3.1

Thermomechanical Processing The thermomechanical processing primarily applied to 9Cr-1Mo grade steel

with ferritic-martensitic microstructure can be broadly categorized into three different temperature domains, Fig. 2.3: Domain-1: Deformation above recrystallization stop temperature (TNR) The

austenite

grains

become

coarse

during

normalization

above

recrystallization stop temperature (TNR) and prior to deformation. These coarse grains get refined in size by several deformation passes above TNR and take part in static recrystallization. Presence of fine microalloying precipitates (such as TiN) particles prevent the grain growth during inter-pass times and maintain fine recrystallized grain structure [15]. After this stage of deformation a dwell time is given during which the temperature drops below TNR. Domain 2: Deformation below recrystallization stop temperature (TNR) Microalloying elements like Nb in steel, retard recrystallisation through staininduced precipitation of fine Nb(C,N) and increases TNR. Deformation passes below TNR (i.e. 950°C-1050°C, depending upon Nb contents in steels [15,50]) produce a heavily deformed austenitic grain structure, with elongated or pancake-shaped grains containing high density of near-planar crystalline defects such as, grain / subgrain boundaries, deformation bands, shear bands and incoherent twin boundaries, which act as the nucleation sites for newly formed finer ferrite grains than by hotrolling [51]. Domain-3: Deformation at temperatures within inter-critical (austenite + ferrite) region Inter-critical rolling is generally applied for obtaining high strength (yield strength and ultimate tensile strength). After deformation in domain-3 regime, a delay time is given to drop the temperature of the steel below Ar3 (austenite to ferrite transformation start temperature). In domain-3, deformation bands continue to form within untransformed austenite grains, which enhance the ferrite nucleation resulting in a finer ferrite grain size [15,50]. The recovery of deformed ferrite grains produces sub- grains within deformed ferrite grains. Smaller sub-grain sizes

12

Chapter 2 increases the strength of the steels, but decreases the ductility and impact toughness due to higher retained strain within the microstructure [15].

Fig. 2.3: Schematic showing the typical rolling schedule and the associated microstructural changes for thermomechanical controlled rolling (TMCR) of steels [50]. 2.3.1.1 Development of crystallographic texture during processing of ferriticmartensitic steel: Different texture components observed in thermomechanically controlled rolled low carbon ferritic-martensitic steels are summarized in Fig. 2.4 [52]. Among the various transformation texture components, shown in Fig. 2.5, the {332} component is reported to be the most beneficial component for achieving good deep drawability and improving strength and toughness. The presence of substitution solutes such as, Mn, Ni, Cr and Mo, finer austenite grain sizes and faster cooling rates during transformation increase the intensity of {332} component. Controlled rolling followed by accelerated cooling, therefore, promotes this texture component. Deformation in ferrite region leads to the formation of alpha (RD||) and gamma fiber (ND||) texture and the recrystallization of deformed ferrite increases gamma (ND||) fiber texture at the expense of alpha (RD||) fiber texture. The austenite recrystallization texture (cube texture, {001}) turns to rotated cube texture component ({001}) in the transformed product. The {001} type of texture has been reported to have an undesirable effect on the delamination behaviour of steels and affect the toughness negatively. But no proper theoretical justification is given in the literature to explain the 13

Chapter 2  

adverse effect of cube texture or the beneficial effect of the gamma fibre texture on toughness. Therefore, detailed understanding on the effect of different crystallographic texture components on toughness is required to develop structural, line-pipe or naval grade steels with excellent combination of strength and impact toughness.

Fig. 2.4: Deformation and transformation textures develop in ferritic-martensitic steel [52].

Fig. 2.5: Different texture components of ferritic-martensitic steel at φ2=45° section of the Euler’s space. 2.3.1.2 Variant selection of martensite during thermomechanical processing: Martensite formation is generally accompanied by a change in volume and a transformation shear. It can, therefore, be anticipated that the course of the 14

Chapter 2 transformation is likely to be affected by the elastic and plastic deformations which result from these phenomena. Depending on the inhomogeneities in temperature, strain, and initial microstructure, as well as on the physical configuration of the specimen, some variants can be favoured over others during transformation. [53–55]. However, if no variants have been preferred during transformation following Kurdjumov-Sachs orientation relationship, the {002} pole figure of product martensite will be as shown in Fig. 2.6. Rules or criteria for variant selection must therefore be postulated, and in order to be successful in predicting the product from the parent texture, models of texture transformation are needed that can also take into account the relative weights of all the variants. Several variant selection models have been proposed over the past few decades. Brief descriptions of all the important ones are given below: 

Shape deformation (SD) model As the martensite formation involves a shear deformation, there is an external

shape change, which has been regarded by Patel and Cohen [56] as an influential factor in variant selection. in the SD model. They considered the interaction of an applied stress with the displacive shear occurring during martensite formation and suggested that the applied stress can aid the total shape change. 

Bokros-Parker (BP) model The BP model suggested by Bokros and Parker [57] is based on the interaction

of slip systems and habit plane variants. During a study of martensite burst formation in strained Fe-Ni single crystals, they observed that the variants that having habit planes perpendicular to the active slip plane were favoured, whilst the variants having habit planes nearly perpendicular to the active slip direction were suppressed. 

Active slip system (AS) model In earlier investigations [58,59] variant selection has been linked to the active

slip systems in the austenite (γ) and in particular to the slip systems subjected to the largest shear stresses during rolling before transformation. The variants produced were those which contained, in the transformation relationship, those slip systems which sustained the maximum resolved shear stress during rolling. 

Twinning shear (TS) model The TS or twinning shear model was first proposed by Higo et al. [60]. They

15

Chapter 2  

proposed that the first shear (i.e. the twinning or twin-forming shear of the γ phase) in the double shear mechanism for the fcc-to-bcc transformation is the deformation which gets affected by the applied stress. This model can account for the observed frequency of the various martensitic variants (α`) produced during stress induced martensite formation by postulating that the critical atomic movement to initiate α` formation is the first shear along the {111}γ direction, followed by a suitable spontaneous complementary second shear that completes the transformation. 

Bain strain (BS) model The Bain strain (BS) model proposed by Furubayashi et al. [61,62] is based on

the interaction between the applied stress (of rolling) and the Bain strain characteristic of the martensitic transformation. They indicated that for the γ to α` transformation, the 'Bain compression axis' (BCA) for each of the three Bain variants is parallel to the [001], [010], or [100] axis of the austenite. It is easy to understand that, if an external compressive stress is applied along the [001] BCA, the variant having [001] as its BCA will be assisted more effectively by the applied stress than the other two. This leads to variant selection. 

Geometrical parameters (GP) model Humbert et al. [63] observed that the rolling reduction or residual stress does

not have a significant effect on the nature of the γ-α` transformation, at least for their samples. They suggested that it was the dimensional parameters of the samples, which influenced the mechanism of variant selection during transformation. They hypothesized two models: in model-1, the selected variants are those which induce maximum deformation along the normal direction of the sheet when the sheet is thin and only contains a few grains across the thickness; in model-2, the selected variants are those which produce the minimum deformation in the plane of the sheet. However, experimental texture should be correlated with the texture predicted from the above mentioned variant selection models, in order to determine the variant selection criteria for a particular process, Fig. 2.7.

16

Chapter 2

Fig. 2.6: Stereographic projection of {002} poles for each of the 24 K-S variants in (111)γ single crystal matrix [63]. The orientation relationships of the symbols and code numbers are shown at right side in tabular form.

Fig. 2.7: (a) Experimental (002) pole figure for transformation texture of martensite obtained from heavily rolled austenite in Fe-25·7Ni alloy [63], and (b-d) simulated (002) pole figures of transformation texture of martensite using the BP, AS, and TS models, respectively [61]. 2.3.2

Normalizing-tempering treatment Thermomechanically controlled rolled plates of high alloy martensitic steel like

9Cr-1Mo steel are usually subjected to normalization treatment to achieve uniform and fine grain structure and to release the rolling strain of the matrix. Typical normalizing treatment comprised of reheating of rolled plate in single phase austenite zone (generally ~ 50 °C above Ae3 temperature) and holding for about an hour for 25 mm thickness, followed by air cooling. In order to achieve fine grain size in the final microstructure, the normalization temperature should be selected in such a way that it is sufficiently high for re-austenitization to take place, but it should not be too high to cause extensive austenite grain growth. The presence of undissolved precipitates at the

17

Chapter 2  

normalization temperature is also beneficial as they restrict the austenite grain growth [15].

2.4

Fracture Mechanism Primarily two types of fracture mechanisms are observed in ferritic-martensitic

grade steels upon impact loading (Fig. 2.8): (i) Ductile fracture and (ii) Brittle transgranular fracture depending upon the test temperature, state of stress and strain rate [64]. 2.4.1

Ductile fracture Generally, ductile fracture is associated with a large amount of plastic

deformation, which results in high energy consumption during this process. Generally, ductile failure occurs by three steps: (i) void initiation, (ii) void growth and (iii) void coalescence. High testing temperature, low strain rate, and plane stress conditions are favorable for ductile fracture. Void initiation happens either by the separation of inclusion-matrix interface (inclusions like, MnS or Al2O3) or by cracking of particles. Voids can certainly nucleate around precipitates like M23C6 and MX in Cr-Mo steel. The shape and/or size of voids depends on the particle size, shape and distribution. The growth phase occurs by the plastic deformation of the matrix surrounding the particle-matrix interface. This stage is the most energy consuming process among the three stages of ductile failure. After a certain amount of void growth, the ligament between the neighbouring voids becomes so small that void coalescence takes place.

Fig. 2.8: (a) cleavage facet forms during brittle transgranular fracture (nucleated from Cr23C6 particle, as indicated by arrow), and (b) ductile fracture, indicated by the presence of micro-voids nucleated from spheroid carbide particles (shown by arrows). 2.4.2

Brittle fracture The brittle or cleavage fracture in Body Centered Cubic (BCC) material

primarily occurs by the transgranular decohesion of {001} cleavage planes [12].

18

Chapter 2 Though the {001} planes are not the low surface energy plane of bcc iron, lower tensile strength along axis due to transformation of BCC to FCC (Face Centered Cubic) during uniaxial tensile loading condition makes {001} plane more vulnerable to cleavage fracture [30]. The favorable conditions for brittle fracture are high strain rate, low temperature, and plane-strain condition. Owing to the asymmetric stress field around the dislocation, the glide stress, i.e. the Peierls-Nabarro stress is higher in BCC materials as compared to the FCC materials [12]. On the other hand BCC materials have only twelve primary slip systems. Hence, the contribution of the glide stress is higher in BCC material as compared to the dislocation interaction stress on the yield strength. As, the Peierls- Nabarro stress is dependent on the temperature, the yield strength of BCC material is also highly sensitive to the temperature, Fig. 2.9.

Fig. 2.9: Schematic diagram showing variation in yield stress and fracture stress with temperature. TGY denotes the general yield temperature. The variations in yield stress and fracture stress as function of temperature are shown in Fig. 2.9. The temperature (TGY), where yield stress intersects with the fracture stress is known as the general yield temperature [64]. The significance of this temperature is that below this temperature the brittle fracture occurs without yielding. Therefore, due to the high temperature sensitivity of yield strength in BCC material, below TGY temperature material first encounter cleavage fracture stress prior to yield stress, and participate in brittle fracture without sufficient yielding. However, the cleavage fracture stress is also a critical parameter, and it depends on the microstructural parameters and micromechanism of cleavage fracture [64].  Basic micromechanisms for cleavage fracture: The process of cleavage fracture in steels involves three critical steps (Fig. 2.10) as stated below:

19

Chapter 2  

(i)

In the first step, a microcrack nucleates at some microstructural inhomogeneities, such as, large inclusions (such as, MnS), harder constituents (like, TiN etc.), grain boundary carbides (such as: Cr23C6), dislocation pile up and twins.

(ii)

Next step is the propagation of this microcrack across the particle-matrix interface.

(iii) In the final stage, the microcrack propagates across the first matrix-matrix boundary. The conditions that dominate each of these processes typically vary with microstructure, temperature, stress and strain distribution ahead of a notch/pre-cracked tip [65,66].

Fig. 2.10: Schematic diagram showing different steps involved in cleavage fracture: Step 1: nucleation of sharp microcrack nucleates at some microstructural feature, Step 2: propagation of microcrack across particle-matrix interface, Step 3: propagation of microcrack across matrix-matrix boundaries [67].

2.5

Instrumented Charpy impact testing Charpy impact testing has been practiced in industry as a common method to

determine the impact transition behavior of steel as function of testing temperature. One of the focuses of the present study is to understand the role of microstructure, precipitated particles and texture on the mechanism of impact fracture properties of low carbon lath martensitic at different temperatures. Typical impact transition curves observed in ferritic-martensitic grade steels in terms of impact energy and fracture appearance are presented in Fig. 2.11 and the characteristic points on the curves are mentioned [49].

20

Chapter 2

Fig. 2.11: Characteristic points on the Charpy impact transition curves. Solid line represents impact energy transition curve and dotted line represents fracture appearance transition curve. USE: Upper shelf energy (USE) represents the ability of a material to absorb fracture energy in complete ductile fracture mode usually observed at room temperature or above. LSE: Lower shelf energy (LSE) represents the absorption of fracture energy in complete cleavage fracture usually observed at cryogenic temperatures (say, 100 °C to -196°C). DBTT: Ductile-to-brittle transition temperature (DBTT) represents the temperature corresponding to the impact energy absorption that is halfway between USE and LSE. FATT: The fracture appearance transition temperature (FATT) represents the temperature at which the fracture surface appearance exhibits 50% fibrous and 50% cleavage fracture. 27J-ITT: 27 J impact transition temperature (27J-ITT) represents the temperature corresponding to the absorbed impact energy of 27 J. NDT: Nil ductility temperature is the temperature below which fracture appearance shows complete cleavage fracture. The load-time or load-displacement curves recorded during the instrumented Charpy impact testing at a particular temperature, show distinct profiles, Fig. 2.12, with characteristic points as discussed below [49]:

21

Chapter 2  

Fig. 2.12: Characteristic points on schematic load-time profile in the transition region. General Yield Load (PGY): The load at which local yielding spreads over the entire area corresponding to the ligament under the notch. Maximum Load (Pmax): It represents the maximum attained load, at which crack extension just takes place by the coalescence of the newly formed voids with pre- existing crack tip/notch. Brittle Crack Initiation Load (Pi): In the transition temperature range, stable crack growth is interrupted at this load and the crack subsequently propagates into the ferrite matrix by stress-controlled cleavage, initiated within the ’trigger‘ particles such as, grain boundary carbides, inclusions and any other ’weak‘ constituents. Crack Arrest Load (Pa): At this point, brittle crack propagation is arrested and reconverted to ductile crack growth under plane-stress condition. The ductile crack further propagates by the formation of shear lips along the slant planes at ~ 45° angle to the macroscopic fracture surface. Load signal recorded in the instrumented Charpy impact testing is superimposed by a scatter oscillation which occurs due to the inertia effect of the hitting hammer. Therefore, it is necessary to subtract the scatter oscillation from the recorded load in order to accurately identify the yield load and the maximum load. Kobayashi has proposed moving point average method to correct the recorded loadtime data [68]:

1 i  m 1 Yi  Xj m  i m

Eqn. 2.1

22

Chapter 2 where, Yi is the smoothed data set corresponding to the sampled data Xj: The smoothing can be performed using values of ‘m’ ranging from 1 to 999 [68,69].

2.6

Effect of microstructure on fracture behavior along the impact

transition regime The effect of temperature on fracture behavior can be summarized in a stress/temperature diagram, Fig. 2.13 [70], which contains different parameters that are dependent on the microstructure. These include the temperature dependency of the yield strength (σ0), the peak stress in the plastic zone (σˆ), as well as, the particle strength (σpm) and matrix strength (σmm). The temperature region over which ductile and cleavage fracture occurs can be clearly identified from this diagram and explained in details below.

Fig. 2.13: Schematic diagram showing the different fracture mechanisms operating at different temperatures determined considering the variation in yield strength (σ0), the peak stress in the plastic zone ((σˆ) and the cleavage fracture stresses considering particles (σpm) and grains (σmm) [70]. (i)

At the lowest temperature ( T3), ductile fracture occurs as the peak stress (σˆ) becomes lower than both the cleavage fracture strength of grain (σmm) and particle (σpm). The voids can generate by the particle cracking at temperature T3, the voids generate by the decohesion of particles as the fracture of particle is not possible at that temperature.

23

Chapter 2  

As the fracture stress of grain-grain or grain-particle depends on grain size inversely following the equation:

 im  

2 E im (1  ) 2  d i

Eqn. 2.2

where, i can be either p (particle) or m (matrix), when referring to particle-matrix interface (p-m) and matrix-matrix interface (m-m), respectively, di corresponds to either particle size, dP, or grain size, dm, and γpm and γmm are the corresponding effective surface energies. For a penny shaped micro-crack, constant

   = 1.25.

Therefore, fracture stress can be enhanced to by reducing the grain size, so that grains starts yielding before it reaches fracture stress. The effect of grain size on fracture stress in ferritic steel was shown in Fig. 2.14, where cleavage fracture stress, σF (earlier mentioned as σmm) is in the range of ~500-1300 MPa for annealed or normalized mild steels (C ~0.03-0.07) for different sizes of grain (~30-800 µm) [71]. The results showed that σF increases with the decrease in ferrite grain size. In case of ferritic-pearlitic structure in TMCR microalloyed steels, it has been found that the average ferrite grain size within range of 5-10 µm results in σF values within in the range of ~1500-2400 MPa [17,72]. Mintz et al. [73] have reported that grain size has a more dominant effect on 27J-ITT then carbide particle size for low C (~0.06 wt.% C) steels. For a certain range of carbide thickness (0.5-1.0 µm) the 27J-ITT was found to drop by ~160ºC, with a decrease in ferrite grain size from ~600 µm to 5 µm, Fig. 2.15. Microalloying along with TMCR not only improves strength but also improves fracture toughness of steel by decreasing the ferrite grain size. Increase in thickness of grain boundary carbides decreases the σF and increases the impact transition temperature, Fig. 2.15.

Fig. 2.14: Variation in cleavage fracture stress (σF, MPa) with D-1/2 showing the increase in σF with the decrease in grain size, D (mm) [71].

24

Chapter 2

Fig. 2.15: Influence of grain size (mm-1/2) and carbide thickness (µm) on 27 J-ITT [73]. Lot of works were done to develop quantitative relationships by regression analysis to predict different impact transition temperatures (ITT), such as DBTT, 27JITT and 54J-ITT for steels considering the influence of factors such as, ferrite grain size (D), carbide thickness (t), pearlite content (%-pearlite), precipitation hardening effects of alloying elements (ΔY), cleanliness, cooling-rate and embrittlement effects. A list of those studies with the empirical equations is given below. 1) Gladman and Pickering proposed a regression equation for the prediction of DBTT (in ºC), which is valid in the composition range 0-0.20 wt.% C, 0-0.013 wt.% N and more than 0.2 wt.% Si [15].

DBTT  19  44(wt %Si)  700(wt %N f )1/ 2  2.2(% pearlite) 11.5D1/ 2 Eqn. 2.3 where, D is grain size in mm, Nf is the free nitrogen content and %pearlite is the pearlite content (up to 30 % pearlite) in ferrite-pearlite steel. 2) Mintz and his co-authors [73] developed relationships between ITT values (27JITT, DBTT, 54J-ITT, and FATT) and microstructural parameters such as grain diameter, D (in mm), pearlite content, and grain boundary carbide thickness, t (in µm), for microalloyed steels. The relationships are:

27J  ITT(ºC) 173t1/2 8.3D1/2  0.37(Y)  42

Eqn. 2.4

54J  ITT(ºC) 161t1/2 11.7D1/2  0.47(Y) 18(%pearlite)1/3  40

Eqn. 2.5

DBTT(ºC) 112t1/2 13.7D1/2 15(%pearlite)0.33  0.43(Y)  20

Eqn. 2.6

50% FATT(ºC) 131t1/2 12.7D1/2  0.45(Y)  46

Eqn. 2.7

where, ΔY is the precipitation hardening component in MPa. 25

Chapter 2  

3) Petch [74] predicted the ITT for ferrite-pearlite steels assuming that cleavage occurs during yielding at a temperature for which the constrained yield stress (YS) equals the cleavage strength (σF). Petch assumed that the strain rate in Charpy test to be ~103 s-1 and predicted the temperature (T, ºC), and grain size dependence of yield strength by following the equation originally proposed by Harding [74].

YS  350  2.5T  K y (D)1/ 2

Eqn. 2.8

Based on eqn. 2.8 the following relationship is developed to predict the 27J-ITT:

5.5(27 J  TT ) (ºC)  770 _ 2.2K y ( D1/ 2 )   F

Eqn. 2.9

where, K y = 21 MPa-mm1/2 Although these equations to relate different impact transition temperatures (ITT) with the ferrite grain sizes are quite established for different ferritic grade steel, nothing of these kinds of correlation exists for materials with martensitic structures owing to difficulty in establishing ‘effective grain’ among the martensitic microstructural units of different length scale (i.e. prior-austenite grain, martensitic packets, martensitic blocks and laths). Therefore, it is necessary first to identify the microstructural unit within hierarchical martensitic grain structures which may act as an ‘effective grain’ for impact fracture resistance.

2.7

Effect of crystallographic texture on impact transition

behaviour High strain rate of deformation associated with impact loading on martensitic steel leads to transition of nature of fracture from ductile to transgranular brittle cleavage fracture mode over a range of temperatures, i.e. known as transition temperature regime. By increasing the fracture stress of any BCC materials, this transition temperature (i.e. TGY in Fig. 2.13) can be further lowered. Therefore, cleavage fracture stress is an important factor for the resistance to brittle fracture in a material. Now, cleavage fracture stress is also dependent on two important parameters, i.e. effective grain size and plastic constraint factor [12], which in turn closely related with the crystallographic nature of the microstructure as discussed below:  Effective grain size Owing to the strong micro-texture developed after thermomechanical processing

26

Chapter 2 of steels number of small grains in the microstructure join together and behave like a single ‘effective-grain’, i.e. a coherent length over which cleavage crack propagates without any deflection along its path [27]. The cleavage crack passes through the lowangle boundaries without any deflection and deflects only when the crack interacts with a high angle boundary. Due to this reason the effective grain size in bainitic / martensitic steels (accelerated cooled / direct quenched / quenched and tempered steels) is determined by the packet size and not by the sizes of individual laths [15– 20]. Acicular ferrite structures show comparatively lower effective grain sizes than the other microstructures and therefore, provided excellent low temperature toughness [75,76]. On the other hand, the effective grain size of dual phase steels can be higher than that of ferrite-pearlite HSLA steels of similar chemistry. Kim and Nakagawa [77] studied the effective grain size in dual-phase ferrite-martensite steels and concluded that the introduction of second phase (martensite) into the ferrite matrix has no effect on changing the direction of cleavage crack propagation; therefore, the effective grain size of dual-phase steels is the same as that of the initial structure prior to the twophase annealing. Kim et al. [24] has demonstrated that effective grain size of a microstructure can be determined from the Electron backscatter diffraction analysis (EBSD) by considering misorientation angle at the grain boundaries. Therefore, determination of ‘effective grain size’ mostly done by considering the grain boundary misorientation angle between the neighbouring crystals [17,18,25–27]. Hence, the effective grain size has been reported to increase with the increase in grain boundary misorientation threshold [17,24]. A number of threshold boundary values in the range of 5° to 25° have been suggested in the literature to separate low-angle and high-angle boundaries [16,24,25,78–80]. A list of different approaches followed in the literature to define the crystallographic grain size in steel is summarized in Table 2.2. Table 2.2: Threshold criterion used for the determination of effective grain size for different microstructures in steel. Threshold criterion on effective grain size

Microstructure

Reference

5° on Misorientation angle

Ferrite-Pearlite,FerriteMartensite

[25]

12° on Misorientation angle

Ferrite-Pearlite

[29]

27

Chapter 2  

15-20° on Misorientation angle

Ferrite-Pearlite, FerriteMartensite, Bainite, Martensite

Largest grain size among Ferrite-Pearlite the grains having {001} plane parallel to fracture surface. Average grain sizes or Bimodal ferrite grain structures largest grain sizes corresponding to the coarseand fine-grain populations.

[24,79,80] [81]

[16]

In order to determine the ‘ effective grain size’, Bhattacharjee et al. [17,29] have carried out EBSD scan on the cleavage fracture surface of the broken Charpy samples, Fig. 2.16. They have demonstrated that a single facet can comprise of more than one grain having less than 12° misorientation between them. In support, they have theoretically shown that a crack loose only 5% of its energy it if deviates by 12°. Inverse pole figure map has also been shown to confirm that the facets formed along the {001} planes of the crystals. Finally it has been concluded that grain boundary misorientation threshold should be set at 12° angle, to estimate the effective grain size.

Fig. 2.16: (a) A cleavage facet from TMCR steel on which orientation imaging (OIM) was carried out; (b) Corresponding OIM image of one cleavage facet from TMCR steel (four different grains are marked); (c) Inverse pole figure of the orientation of crystals detected on fracture surface [17]. However, in case of earlier investigations on martensitic steel, Hughes et al. [21] indicated cleavage crack propagation across prior-austenite grain boundary and lath boundaries, whilst others [22–24] reported prior-austenite grain boundaries as

28

Chapter 2 effective in crack retardation. Most of the earlier research showed martensitic packet boundaries as effective barrier to cleavage crack propagation [25–28]. These investigations either correlated the size of cleavage facets in transgranular fracture modes with the packet size and block size [19,21–23] or estimated the ‘effective grain’ size considering only the high-angle boundary [17,18,25–27]. In these literatures, various types of boundaries (e.g. packet boundaries and block boundaries) are differentiated in terms of angle of misorientation (angle-axis) exist in those boundaries, which may be misleading for characterizing their effectiveness in cleavage crack retardation as cleavage crack passes through {001} type of planes in neighbouring crystals for body-centred cubic (BCC) materials [24,25,29,30]. Moreover, the role of block boundaries are still unclear from this regard. However, Guo et al. [31] considered the theoretical aspects of Bain variants and their possible effect on cleavage fracture. As per the present knowledge of the authors, there are very few studies available that dealt with direct observation of cleavage crack path in view of crystallography of martensite [27,32].  Plastic constraint factor (PCF) The factor, KPCF (PCF), which influence the cleavage fracture stress of a material is the ratio of maximum principle stress ahead of the crack/notch to the yield stress [82]. Different values of PCF have been reported in literature as listed in Table 2.3. Table 2.3: Different Plastic Constraint Factor (PCF) values reported in the literature. Authors (Reference)

KPCF

Ewing et al. [83]

1.94

Shindo and Tomita [83]

1.80

Green and Hundy [83]

2.18

Chen et al. [84]

2.52

Moitra et al. [85]

2.57

Although, the PCF depends on the local crystallographic orientation of the material ahead of the crack tip, the effect of crystallographic orientation on the plastic constraint factor has not been considered earlier. However, in a recent work it has been found that local crystallographic orientation or in other words, crystallographic texture

29

Chapter 2  

strongly influence the plastic constraint factor [86]. Therefore, fracture stress is dependent on the texture of the microstructure, which eventually determines the nature of impact fracture (brittle or ductile) and provides a direction for improvement of impact fracture resistance. Ghosh et al. in their model considered a crystal with certain orientation ahead of the crack tip, Fig. 2.17. The angle between the cleavage planes 1, 2 and 3 (as marked on Fig. 2.17) and the fracture planes are h1, h2 and h3, respectively. It was considered that among these three angles, h1 is the lowest. The cleavage plane of the crystal (say, plane 1), which makes the minimum angle (hmin = h1) with the fracture plane of the sample, is considered to be the active cleavage plane. Hence, the cleavage crack will actually propagate through cleavage plane 1. Considering angle hmin, the plastic constraint factor (PCF) was evaluated using the following equation:

PCF  (1 n  m mn  m2  n2 )1/2

Eqn. 2.10

where, m  (1  sin  min ) /(1  sin  min ) and n  2 /(1  sin  min ) 2

2

2

And m is the Poisson’s ratio. The angle hmin will vary depending on the orientation (G) of the crystals in front of the notch. Therefore, the plastic constraint factor will be different for different crystal orientations.

Fig. 2.17: Schematic diagram showing a crystal with an arbitrary orientation with respect to the fracture plane of the sample. RD, TD and ND are the rolling direction, transverse direction and normal direction, respectively.[86].

2.8

A review on studies on impact transition behaviour of 9-12 wt%

Cr-Mo steel As out lined in Section 2.2, the primary challenge for the application of 912%Cr-Mo steel is its nature of transition from ductile fracture mode to brittle transgranular fracture mode when this steel experiences irradiation, Fig. 2.2. In this

30

Chapter 2 section, a short review on various works done to investigate the ductile-to-brittle transition behaviour of 9-12%Cr-Mo steel using Charpy impact tests is addressed. Using small ball punch test Misawa et al. [87] showed that ductile-brittle fracture energy transition behavior can be determined, and correlated small punch test DBTT with the DBTT measured from Charpy-V notch tests. Yoshida et al. [88] studied U-notch half-size using instrumented Charpy test on a 9Cr-1Mo and a 9Cr-2Mo ferritic steels. The irradiated specimens showed a shifts in DBTT for 30 to 40 K. Kayano et al [89] studied the embrittlement phenomena on different Fe-Cr ferritic steels due to neutron irradiation with the help of instrumented Charpy and tensile tests, and established equations were to correlate strength, ductility, DBTT and fracture toughness (J values). Vitek et al. [89] indicated that the largest effect on the Charpy impact properties of a 12Cr-1MoVW alloy were found after irradiation at temperature between 300°C and 400°C. However, Klueh and Alexander [90] found higher increase of DBTT and decrease of USE for irradiation at 400°C as compared to at 300°C. The shifts of the DBTT were of the order of 204°C for a V, Nb added P91 steel. Moreover in a review Klueh et al. [91] charted that the effect of irradiation can be controlled by the initial heat-treatment, and the irradiation-induced shift of the DBTT can be reduced by addition of Tantalum that promotes refinement of microstructure. Kayano et al. [92] studied the embrittlement effect with Charpy tests on Fe-Cr ferritic steel, low activation steels based on Fe-Cr-W composition, and Fe-Cr-Mo, Nb or V ferritic steels, and found lower irradiation-induced shift of the DBTT in the composition range 3-9% Cr. Klueh and Alexander [93] also observed lowest shift of DBTT in 9%Cr steels among 9-12% Cr grade steels. Odette et al. [94] ran a series of finite element simulation with combination of the crack tip stress field calculations and weakest link statistics to model cleavage fracture in ferritic-martensitic grade steel. Odette et al. also developed a local approach based micro-mechanical model for quasi-cleavage in tempered martensitic steel involving a critical stress σ* surrounding a critical area A* around the crack tip [47]. They found that the underlying mechanism is a stable- to-unstable transition of a process zone formed by the coalescence of many micro-cracks that arrested in the tempered martensitic microstructure. To study the effect of the loading mode on the critical J-integrals, Li et al [95] developed a modified compact tension specimen geometry for Cr-Mo steels. Edsinger et al. [96] highlighted the fact that depending on

31

Chapter 2  

the specimen size or the deepness of the crack and the loading rate, different transition temperatures can be found for martensitic steel. Odette and co-authors further continued their work on developing models to account for the influence of the specimen size and geometry on fracture [97]. In a recent work, Odette et al. [98] emphasised to relate the critical stress σ* to both the intrinsic arrest fracture toughness of the matrix and an effective brittle particle size. Spätig et al. [99] compared the previously published data with the data they reported, and proposed a temperature adjusted master curve approach. In different international research program, several works were done to measure fracture toughness [48,100,101] with a variety of specimen sizes, geometries and testing conditions based on the philosophy of the ASTM E1921 standard, and results indicated an apparent high scatter in the database. To follow the ASTM E1921 standard method for determination of fracture toughness is quite time consuming, on the other hand relatively quick method to assess the fracture characteristic of any material is determination of DBTT and upper-shelf energy using Charpy impact test. Lechtenberg [102] mentioned that the fracture properties of ferritic steels for fusion applications were generally determined using measurement of DBTT by Charpy tests, but there is no quantitative method to use the Charpy data in design. Therefore, a method should be developed which would be quick as that of Charpy impact test, and at the same time that data can be used for design purpose during fusion reactor application. Hosoi et al. [102] has undertaken a study to investigate the effect of long-term aging on the microstructure and toughness of different Cr-2Mo steel using V-notched instrumented Charpy impact testing. They observed a large loss in toughness values associated with the precipitation of Si and P containing intermetallic compounds in the ferrite. Harrelson et al. [103] also found that these impurity elements (S, P and Si) increase the DBTT and decrease the upper shelf energy in 9Cr steels. However, Vitek et al. [104] assessed the aging effect on the DBTT and upper shelf energy of 9-12%CrMo steel and observed that these properties are not affected for samples that are aged at 300°C for 2500 h., Moreover, Clausing et al. [105] studied the effect of chemical segregation of Ni, Cr, Si and P on fracture behaviour of martensitic steels. They found that the segregation of these elements along the prior-austenite grain (PAG) boundaries significantly reduced the fracture resistance. Using different austenitization temperature Maiti et al. [106] changed the PAG size and martensitic packet size, but

32

Chapter 2 found that the lower shelf fracture toughness was quite insensitive to the changes in the PAG size and martensitic packet sizes. In order to model the role of carbides in brittle fracture in martensitic steel, Lucas et al. [107] considered a critical stress to propagate a carbide-size crack into the surrounding matrix. They employed different kinds of heat treatments to vary the prior austenite grain sizes, packet sizes, and carbide distributions and densities and found that cleavage fracture is facilitated by microcrack nucleation and propagation from the large boundary carbides. It was observed that the lowest DBTT and highest upper shelf energy was achieved at tempering temperature of 750 °C compared to 650 °C for 9Cr-1Mo steel [108]. Moreover, optimisation was done by varying time and temperature of the austenitising and tempering treatments to refine the prior austenite grain size and reduce the strength of the matrix of the 9Cr steels in order to have better impact properties [109]. As the cleavage crack can nucleate easily from any of the numerous carbide particles, the DBTT of ferritic-martensitic steels primarily depends on the resistance offered by the microstructural barriers, i.e. the grain boundaries, to the crack propagation [7]. Now, martensite in 9Cr-1Mo steel has a hierarchical microstructure comprised of different structural units of varying length scales like lath, sub-block, block, packet and prior-austenite grains. The boundaries separate these microstructural units are different in nature and have different misorientation angles across them [10,11]. Earlier studies showed the beneficial effect of refining the prior-austenite grain size (PAGS) and martensitic packet size (MPS) in lowering the DBTT as their boundaries effectively retard / divert the cleavage crack propagation [16–18]. On the other hand, martensitic lath boundaries are known to be ineffective from this respect [27]. The role of sub-block and block boundaries on DBTT is not yet established. Besides that ‘whether all the packet boundaries are beneficial or not?’ is yet to be understood. Therefore, impact properties (DBTT and upper-shelf energy) of 9Cr-1Mo steel should be correlated with different types of martensitic grains (i.e. prior-austenite grain, and martensitic packet, block and laths) to establish a proper relation, which can further be used to improve the impact properties by tailoring that ‘effective grain’ size.

2.9

Creep Mechanism Creep is a time dependent plastic deformation that a material experiences at a

constant load or stress. This phenomenon usually occurs in a material at temperatures higher than room temperature, unless room temperatures are high homologous 33

Chapter 2  

temperature for that material. Generally creep is represented by strain-time curve, as shown in Fig. 2.18.

Fig. 2.18: A typical creep curve indicates three different regions: the primary, secondary and the tertiary creep region [49]. Once the material experiences an instantaneous strain, ɛ0, which is composed of elastic (recoverable on release of load), anelastic (recovers with time) and plastic (nonrecoverable) strain, as a result of sudden loading, the primary creep region begins only after that. As the name suggests primary creep region describes the initial stage of creep deformation. This region is characterized by a decreasing strain rate with time. This continues untill the secondary stage starts. The strain rate of deformation remains constant during the secondary creep region. The strain rate in secondary stage is the minimum strain rate of a creep deformation. The creep life of any material can be estimated through the knowledge of the creep strain rate in secondary stage. The last stage of creep deformation is the tertiary creep regime. In this stage, the material undergoes very high strain rate deformation and eventually fractures. The substantial amount of the total strain experienced by a material is contributed by the secondary stage creep regime. The creep curve is a result of microstructural level changes occurring in a material. This curve indicates a competition between the processes of strain hardening and recovery. Materials usually get strain hardened during plastic deformation. For further plastic deformation the applied stress must exceeds the increase in flow stress of the material due to strain hardening. Otherwise, deformation can also proceed at the initial applied stress if the material softens. The recovery process acts as a softening mechanism in a deformed specimen to allow further plastic deformation. Hence, creep in a material is the consequence of a competition between the mechanisms of strain hardening and recovery [12].

34

Chapter 2 2.9.1

Standard Creep equations The time dependent creep rate can be described by [110]:

   aiti ni

Eqn. 2.11

i

where,  is the creep rate, ai and ni are functions of both temperature and stress. The primary stage of a creep curve can be described by eqn. 2.11 when n reaches a value of 2/3. In such a situation, the creep strain (ε) as function of time can be described as:

  0  t1/3

Eqn. 2.12

where, ε0 is the instantaneous strain, β is a constant and t is time. Eqn. 2.12 is in agreement with the time law of creep proposed by Andrade [111]. However, this equation is only valid for higher temperature creep, where recovery is significant. For steady state creep, n = 0. The creep curve can be then described as:

 0 St

Eqn. 2.13

where, εS is the steady state creep rate. Therefore, the total creep strain, combination of the primary, secondary and tertiary region can be described as:

  0  t1/3  St   t3

Eqn. 2.14

where,  t describes the tertiary component of the creep curve. 3

From the above discussions, it is quite evident that creep deformation is a function of the applied stress, temperature and the initial microstructure of the material. Consequently, the strain rate of deformation,  can be expressed as:

 = f (σ, ,T, microstructure)

Eqn. 2.15

where σ is the applied stress and T is the test temperature. 2.9.1.1 Effect of stress and temperature The steady state strain rate of creep deformation has been found to be directly dependent on the applied stress for a given temperature. This relation can be expressed by Norton’s law [112]:

S  K n

Eqn. 2.16

where K is a constant and n is the stress exponent.

35

Chapter 2  

Similarly, the rate of creep deformation increases with the increase in temperature for a constant applied stress. The effect of temperature on creep rate is indicated as following equation: S  K  n exp(

QC ) RT

Eqn. 2.17

where K′ is another constant and QC is the activation energy of deformation. The activation energy is dependent upon the mechanism that controls the creep rate. The effect of stress and temperature is represented in Fig. 2.19. With the increase in stress and temperature, the instantaneous strain at the time of stress application increases, the steady state creep rate is increased and the time to rupture (t) is also diminished.

Fig. 2.19: Illustration of the effect of stress and temperature on creep behavior of any Material. 2.9.1.2 Effect of microstructure In general, the analysis of any creep data is done by assuming the microstructure to be constant. Some of the microstructural features that are prone to change during the course of creep deformation are precipitate size and distribution, composition of phase, and grain size. Hence, in order to estimate the different creep parameters and the mechanism of creep, it is necessary to consider the microstructure to be constant. Even though thermal stabilization establishes a constant microstructure during the course of a test, as the materials are usually heat treated at temperatures higher than the test temperature, stress assisted processes altering the microstructure cannot be ruled, Fig. 2.20. Moreover, in case of non-equilibrium structures such as nanocrystalline materials undergo stress assisted microstructural changes, which prevent these materials to attain a constant creep microstructure. Therefore, it is recommended that creep tests on nanocrystalline material should be carried out at a stress level that is lower than the critical stress at which microstructural changes could be initiated.

36

Chapter 2

Fig. 2.20: Coarsening of lath width after long term creep exposure in a ferriticmartensitic steel. (a): Virgin material, and (b) material in crept condition [113]. The most important microstructural parameter that generally plays a major role in controlling the creep properties is the grain size. Although the Hall-Petch relation states that the materials with finer grain sizes possess greater strength than materials with larger grain size, under creep conditions the reverse is true. Finer grain sized materials undergo faster creep damage than coarse grained materials at lower stresses and at higher temperatures. There are certain creep mechanisms which operate faster in materials with smaller grain as compared to materials having coarse grains. The effect of grain size on the steady state creep rate of a material is expressed as follows: S  K d  P n exp(

2.9.2

 QC ) RT

Eqn. 2.18

Identifying the mechanisms of creep With the knowledge of the grain size exponent (p), the stress exponent (n) and

the activation energy (Qc), the mechanism of creep is possible to be identified using the following Table 2.4.: In addition to these three parameters, the mechanism of creep in a material can be identified through the knowledge of the creep constant A from eqn.

2.19. Distinct values of A are found for each creep mechanism.

 kT

P

 b     A    DEb d  E

n

Eqn. 2.19

37

Chapter 2  

Table 2.4: Identification of the particular mechanism of creep through knowledge of the creep parameters, n, p and QC. Where Qgb is the grain boundary diffusion activation energy and Q is the lattice diffusion activation energy

Broadly mechanisms of creep can be classified in two types: (i) diffusion based processes and (ii) Dislocation based processes. Coble and Nabarro-Herring (N-H) are the two deformation mechanisms that fall under the category of diffusion based processes, whilst Harper-Dorn (H-D), Viscous glide and Dislocation climb are following dislocation based mechanisms during creep. In case of Grain Boundary Sliding processes (GBS), combination of diffusion and dislocation based mechanism plays role during creep deformation. However, the mechanism of Power Law Breakdown is not well understood yet [114]. Stress exponent (n) value of 1 suggests that the mechanism of creep deformation could be Coble, Nabarro-Herring or Harper-Dorn. Moreover, with the knowledge of the grain size exponent (p) or the activation energy (QC) the mechanism of creep can be certainly established for any material. For example, a stress exponent of one and an activation energy equal to the lattice diffusion activation energy would suggest the mechanism of creep to be either N-H or H-D. However, if the grain size exponent is equal to 2, then it would establish that the mechanism of deformation is NH. On the other hand, if the steady state strain rate is found to be independent of the grain size (p = 0), then the mechanism of creep is H-D. In addition to these parameters a knowledge of the creep constant A can help in identifying the exact creep mechanism,

Table 2.4.

2.10 A short review on creep behaviour studies of 9-12%Cr-Mo steel Previous studies on creep behaviour of Cr-Mo steels have extensively dealt with the mechanisms of creep deformation depending upon different operating temperature and stress regimes [115–126]. Another aspects of studies have been the effect of grain structure [127,128], stability of different types of carbide (i.e. M23C6 or MX types) particles [113,129–131] and the laves phases evolved during thermal 38

Chapter 2 exposure for long-duration [126,132–134], and special boundaries [135]. Numerous studies on 9Cr-1Mo steel investigates the evolution of microstructure in terms of formation of subgrains and the nature of these boundaries, variation in dislocation density and formation of different types of precipitates [113,136–141]. However, in order to use in ultra-supercritical (USC) power plants or in fuel tubes of fast breeder nuclear reactors, heat resistant modified 9Cr-1Mo steel is processed through various thermo-mechanical

processing

like

heat

treatment,

hot-deformation

etc.

[18,20,28,118,142–146], and finally used in normalized and tempered condition [147]. Normalization and subsequent tempering treatments change the size of the martensitic grain structures, precipitation behavior and crystallographic texture [19,148]. Therefore, it is necessary to study the effect of different normalization treatments on the creep properties of 9Cr-1Mo steel in order to optimize the processing parameter before material is subjected to creep deformation at different temperature regime. In the present investigation, the effect of different normalization temperatures on creep properties and the evolution of microstructures and crystallographic texture after creep tests at different temperatures are studied.

2.11 Summary From the discussions of earlier literatures it has become clear that the microstructural units (like prior-austenite grain, martensitic packet, block and lath), which can act as an ‘effective grain size’ for improving the impact toughness behaviour of martensitic steel is known yet. Therefore, it is required to alter the martensitic microstructure of 9Cr-1Mo steel using hot-rolling and normalization treatments in order to change the dimension of these microstructural units and establish a correlation between impact toughness properties and different martensitic units to identify the effective microstructural unit that will improve the impact fracture behaviour of 9Cr1Mo steel. Moreover, the creep properties of heat resistant 9Cr-1Mo steel should not get much affected by the thermomechanical processing applied to the material in order to achieve better impact toughness properties. Therefore, an optimization of processing should be obtained which would provide better impact toughness properties without much deteriorating creep properties of 9Cr-1Mo steel. In the results and discussions section, these issues have been addressed.

39

 

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40

 

Chapter 3 Experimental details

 

 

Chapter 3  

3.1

Materials A modified 9Cr-1Mo steel has been studied in the present investigation. The

chemical compositions of the investigated steel is given in Table 3.1. The investigated steel is called as ‘modified’ because of addition of microalloying elements like Nb, V and Ti to the conventional 9Cr-1Mo steel [149,150]. The present steel is a low-C (0.1 wt%) steel having 9 wt% Cr and 1 wt% Mo. Microalloying elements such as, Nb, V and Ti is added in this steel to form finely dispersed and stable MX types of carbides or carbonitrides along the martensitic lath boundaries during tempering treatment of air-hardened martensitic structure [149,150]. Table 3.1: Chemical composition of the investigated samples, (wt.%). Element C Wt. %

3.2

Mn

P

S

Si

Cu

Ni

Cr

Mo

Al

Nb

V

Ti

The modified 9Cr-1Mo steel has been received from IGCAR, Kalpakkam in processed by hot-deformation followed by normalized and tempered condition. The hot-rolled material was normalized at 1050ºC for 25 min, air-cooled and tempered at 750ºC for 75 min. Therefore, the ‘As-received’ material has already been austenitized once after rolling. The processing schedules applied on the as-received steel in the present investigation can be divided primarily into four major studies, as are mentioned below: Study on the effect of hot deformation schedule on the microstructure 

Fe

0.1 0.41 0.016 0.005 0.21 0.06 0.20 8.94 0.86 0.01 0.08 0.2 0.003 0.05 Bal

Processing and heat treatment schedules:

3.2.1

N

Small blocks of (20 mm×15 mm×10 mm) were cut from the quarter

thickness of ‘As-received’ plate (400 mm×250 mm×25 mm) and were subjected to reheating [1403K (1130°C) for 3 minutes] followed by plane-strain compression testing at high temperature [1323K (1050°C) - 1148K (875°C)] inside Gleeble®3500 thermo-mechanical simulator. Both single-pass and multi-pass deformation schedules were applied in Gleeble® as schematically shown in Fig. 3.1. In the single-pass schedule, samples were deformed at different deformation temperatures as mentioned in Fig. 3.1a and water-quenched immediately after deformation. 41

Chapter 3  



In multi-pass schedule, three samples (MP1, MP2 and MP3) were tested

and the total deformation applied (Ɛ=0.8) was identical to the single pass schedule. But instead of a single-pass, the total deformation was divided into four (4) identical passes (Ɛ/pass=0.2, Ɛ ̇=1/s) applied at different temperatures, Fig.3.1b. Temperatures of the first pass were within a close range [1313K (1040°C) - 1323K (1050°C)] for all the multi-pass schedules but the subsequent passes were applied at 10 s, 30 s and 50 s intervals (i.e. inter-pass time). Maintaining a constant cooling rate of 1 K/s during sample cooling inside Gleeble®, the deformation temperatures were therefore, separated by 10K, 30K and 50K for the multi-pass schedules and the finish deformation temperatures (FDT) were 1273K (1000°C), 1223K (950°C) and 1173K (900°C), respectively. The samples were water quenched immediately after the finishing passes were applied. In order to understand the sequential microstructural changes taking place during multi-pass deformation, number of samples have been deformed upto different stages of multi-pass deformation (i.e. first, second and third pass) and water quenched, Fig. 3.1b. During each Gleeble schedule, couple of Thermocouples were attached to the sample for carefully monitoring the sample temperature at an accuracy of ± 2 K. Depending on the deformation temperatures, single-pass (SP) deformation schedules were coded as SP-1323, SP-1273, SP-1223 and SP-1173. Depending on the finish deformation temperatures (FDT), the multipass schedules were coded as: MP1-FDT1273, MP2-FDT1223 and MP3-FDT1173. Depending on the total applied strain, the samples quenched from different stages of multi-pass schedules were coded as: MP1-0.2, MP1-0.4, MP1-0.6, MP1-0.8, MP20.2, MP2-0.4, MP2-0.6, MP2-0.8, MP3-0.2, MP3-0.4, MP3-0.6 and MP3-0.8, Fig.

3.1b. Thus the final deformation temperature of MP1-0.4 sample is 1303K (1030°C), MP2-0.4 sample is ~1283K (1010°C), MP3-0.4 sample is ~1273K (1000°C), and so on.

42

Chapter 3  

Fig. 3.1: Schematic diagrams showing (a) single pass deformation schedules, (b) multi pass deformation schedules. 

A small block (20 mm x 15 mm x 10 mm) was cut from an ‘As-

received’ plate and was subjected to 15 pass deformation schedule in plane-strain compression mode inside Gleeble 3500® simulator with true strain of 0.1 per pass starting from 1125ºC down to 775ºC with an interval of 25ºC at the cooling rate of 1ºC/s, Fig. 3.2a. The mean flow stress (MFS) calculated from the stress-strain curve for each deformation pass was plotted against the inverse of temperature and the temperature correspond to the point of change in slope represents the recrystallization stop temperature (TNR) of the investigated steel (determined to be ~ 950ºC), Fig. 3.2b.

Fig. 3.2: (a) Schematic representation of 15 pass deformation schedule applied to the ‘As-received’ steel using Gleeble 3500® thermo-mechanical simulator, and (b) the corresponding mean flow stress (MFS) of each pass plotted against inverse of deformation temperature to determine the point of change in slope, which represents recrystallization stop temperature (TNR). 3.2.2

Study and effect of hot-rolling to vary the microstructure and properties

43

Chapter 3  

The As-received plates were subjected to soaking treatment (1130ºC for 1 h) followed by hot-rolling at four different deformation temperatures: 1050ºC, 1000ºC (both above TNR), 950ºC (around TNR) and 875ºC (below TNR), Fig. 3.3. The deformation was applied in single rolling pass using ~50% strain (true strain of ~0.7), reducing 25 mm As-received plate to 12.5 mm thick rolled plates, over a temperature range (~30ºC) and the finish rolling temperatures are reported above. Depending on the finish rolling temperatures the plates are coded as 1050-HR, 1000-HR, 950-HR and 875-HR. The rolled plates were finally tempered at 750ºC for 1h, Fig. 3.3.

Fig. 3.3: Schematic diagram showing the heat-treatment and rolling schedules used in the present study. 3.2.3

Study on the effect of normalization temperature on the microstructure

and properties The As-received plates were reheated at the heating rate of 10ºC/min to three different austenitization temperatures: 950ºC (950-Reheat), 1025ºC (1025-Reheat) and 1100ºC (1100-Reheat). The heat treated blocks were soaked for 1 h, air-cooled to room temperature and finally tempered again at 750ºC for 1 h. The heat-treatment schedule employed in the present study is summarized in the schematic diagram in Fig. 3.4.

Fig. 3.4: Schematic diagram showing the heat-treatment schedule used is the present study. 44

Chapter 3  

3.2.4. Study on the combined effect of hot-rolling and normalization temperature on the microstructure and properties The as-received modified 9Cr-1Mo steel plates were subjected to soaking (1130 °C for 1 h) followed by controlled hot-rolling (rolling strain ~60%) at four different temperatures: 1050 °C, 1000 °C 950 °C and 875 °C. Each hot-rolled plate was further normalized by austenitizing and air-cooling from three different temperatures: 1100 °C, 1025 °C and 950 °C (for 1 h). Rolled and normalized samples were finally tempered at 750°C for 1 h. A schematic representation of the thermo-mechanical processing is given in Fig. 3.5. Hereafter, these twelve investigated samples (4 hot-rolling×3 normalizing) will be referred using sample codes mentioning hot-rolling temperature (HR) followed by normalizing temperature (NT) as indicated in Fig. 3.5.

Fig. 3.5: Schematic representation of different thermo-mechanical processing routes with different colors being used for samples deformed at different hot-rolling (HR) temperatures. All the sample codes based on their heat treatments and thermo-mechanical processing history are given in Table 3.2

Table 3.2: Sample codes with processing details along with the section numbers where these are mentioned. Section.

3.2.1

Sample code

Sample Details

SP1323

Single pass deformation, strain=0.8, Hot-deformation at 1323K (1050ºC)

SP1273

Single pass deformation, strain=0.8, Hot-deformation at 1273K (1000 ºC)

SP1223

Single pass deformation, strain=0.8, Hot-deformation at 1223K (950 ºC)

SP1173

Single pass deformation, strain=0.8, Hot-deformation at 1173K (900 ºC)

MP1-0.2 / MP2- Strain = 0.2 /pass, total strain = 0.2, deformation temperature 1050 ºC 0.2 / MP3-0.2 MP1-0.4 Strain = 0.2 /pass, total strain = 0.4, deformation temperature 1040 ºC MP1-0.6

Strain = 0.2 /pass, total strain = 0.6, deformation temperature 1030 ºC

45

Chapter 3  

3.2.2

3.2.3

MP1-0.8

Strain = 0.2 /pass, total strain = 0.8, deformation temperature ~1000 ºC

MP2-0.4

Strain = 0.2 /pass, total strain = 0.4, deformation temperature 1020 ºC

MP2-0.6

Strain = 0.2 /pass, total strain = 0.6, deformation temperature 990 ºC

MP2-0.8

Strain = 0.2 /pass, total strain = 0.8, deformation temperature ~950 ºC

MP3-0.4

Strain = 0.2 /pass, total strain = 0.4, deformation temperature 1000 ºC

MP3-0.6

Strain = 0.2 /pass, total strain = 0.6, deformation temperature 950 ºC

MP3-0.8

Strain = 0.2 /pass, total strain = 0.8, deformation temperature ~900 ºC

875-HR

Finish rolled at 875ºC and air cooled, tempered at 750ºC and air cooled

950-HR

Finish rolled at 950ºC and air cooled, tempered at 750ºC and air cooled

1000-HR

Finish rolled at 1000ºC and air cooled, tempered at 750ºC and air cooled

1050-HR

Finish rolled at 1050ºC and air cooled, tempered at 750ºC and air cooled

950-Reheat

Heat treatment at 950ºC and air cooled, tempered at 750ºC and air cooled

1025-Reheat

Heat treatment at 1025ºC and air cooled, tempered at 750ºC and air cooled

1100-Reheat

Heat treatment at 1100ºC and air cooled, tempered at 750ºC and air cooled

875HR-950NT Finish rolled at 875ºC and air cooled, normalized at 950 ºC and air cooled, tempered at 750ºC and air cooled

3.2.4

3.3

875HR-1025NT Finish rolled at 875ºC and air cooled, normalized at 1025 ºC and air cooled, tempered at 750ºC and air cooled 875HR-1100NT Finish rolled at 875ºC and air cooled, normalized at 1100 ºC and air cooled, tempered at 750ºC and air cooled 950HR-950NT Finish rolled at 950ºC and air cooled, normalized at 950 ºC and air cooled, tempered at 750ºC and air cooled 950HR-1025NT Finish rolled at 950ºC and air cooled, normalized at 1025 ºC and air cooled, tempered at 750ºC and air cooled 950HR-1100NT Finish rolled at 950ºC and air cooled, normalized at 1100 ºC and air cooled, tempered at 750ºC and air cooled 1000HR-950NT Finish rolled at 1000ºC and air cooled, normalized at 950 ºC and air cooled, tempered at 750ºC and air cooled 1000HR-1025NT Finish rolled at 1000ºC and air cooled, normalized at 1025 ºC and air cooled, tempered at 750ºC and air cooled 100HR-1100NT Finish rolled at 1000ºC and air cooled, normalized at 1100 ºC and air cooled, tempered at 750ºC and air cooled 1050HR-950NT Finish rolled at 1050ºC and air cooled, normalized at 950 ºC and air cooled, tempered at 750ºC and air cooled 1050HR-1025NT Finish rolled at 1050ºC and air cooled, normalized at 1025 ºC and air cooled, tempered at 750ºC and air cooled 1050HR-1100NT Finish rolled at 1050ºC and air cooled, normalized at 1100 ºC and air cooled, tempered at 750ºC and air cooled

Microstructural characterization Microstructural characterization of the investigated samples has been carried

out by optical microscopy, scanning electron microscopy (SEM), electron backscatter

46

Chapter 3  

diffraction (EBSD) analysis and transmission electron microscopy (TEM). Samples collected from RD-TD and RD-ND planes of the rolling plates (Fig 3.6), were prepared for optical microscopy and image analysis by mounting on Bakelite followed by grinding and polishing down to 0.25 µm Al2O3 finish. The samples were etched with either Vilella’s reagent (100 mL ethanol + 5 mL HCL+1 g picric acid) or hot- and saturated-picric acid solution for revealing the martensitic microstructure and the prior-austenite grain boundaries, respectively. Microstructural study has been carried out using Leica ®DM6000M optical Microscope, fitted with Leica® M.W. and Leica® L.A.S image analysis software and using Zeiss® EVO60 and Zeiss® Auriga Compact Dual-beam Scanning Electron Microscopes (SEM) fitted with Oxford-Inca® PENTA FETX3 for Energy Dispersive X-ray Spectroscopy (EDS). The average of prior-austenite grain size (PAGS) and martensitic packet size were measured through equivalent circle diameter (ECD) method based on the determination of grain area (by image analysis) of more than 500 grains from each sample following the procedure reported by Chakrabarti et al. [151]. The volume fractions of second phases like precipitates (M23C6) and inclusions (MnS or Al2O3) were also estimated using the image analysis technique by considering at least 500 µm × 500 µm area for each sample Samples collected from RD-TD and RD-ND planes of the rolling plates (Fig 3.6) were subjected to electron backscatter diffraction (EBSD) analysis. Sample preparation for EBSD consists of standard mechanical polishing (down to 0.25-μm grit), followed by electro-polishing with the solution containing 5 vol.% perchloric acid and 95 vol.% acetic acid. The electro-polishing has been carried out in Buehler Electromet4® for 15 s at 20 Volt. EBSD studies were also carried out on the fracture surfaces of broken Charpy specimens to study the cleavage facets and on the transverse section, perpendicular to the fracture surfaces, to study the secondary cracks present just beneath the fracture surfaces. Sample preparation was not required for the EBSD study on cleavage facets. The EBSD scanning has been performed using HKL Channel 5 system (Oxford Instruments, Abingdon, Oxfordshire, UK) attached to Zeiss® Auriga Compact SEM at 70° tilting condition. At least 500 µm × 500 µm area was scanned for each samples at different step sizes ranging between 0.1 µm and 0.3 µm, depending upon the requirement. HKL Channel 5 software and MTex toolbox were used for the EBSD data analysis. 47

Chapter 3  

Fig. 3.6: Typical dimension and orientation of samples for microstructural study, EBSD study, texture study, tensile testing and Charpy impact testing with respect to the original rolled plate

3.4

Texture study The macro-texture study was conducted on the RD-TD plane of the

samples, taken from the ¼ thickness location of the rolled plates, Fig 3.6. Samples were further polished down to 0.25 µm Al2O3 finish before the scanning. The XRay scanning has been carried out at 5° step size for a range of Φ=0°-360° (rotation about ND axis) and ψ=0°-80° (rotation about TD axis) for four crystallographic planes of each sample using PANanalytical® X-ray goniometer. Co Kα radiation has been used for the diffraction. MTex toolbox has been used to analysis the collected pole figures. Background and defocusing correction have been made before the construction of orientation distribution function (ODF).

3.5

Hardness testing Hardness values were determined on the plane parallel to the RD-TD rolling

plane at 1/4 th thickness location. In order to get the accurate measurement, both lower and upper surfaces were flattened. The specimens were polished down to 0.25 µm Al2O3 finish and etched with 2%-nital solution before the testing. Macro-hardness was measured at 20 kgf load using a LV-700 model LECO Vickers tester. The macrohardness value for each sample was determined based on the average of twenty five (25) readings.

3.6

Uniaxial tensile testing and tensile creep testing Tensile specimens for uniaxial tensile as well as for tensile creep tests were

48

Chapter 3  

prepared along the longitudinal orientation (i.e. tensile axis // RD) from the samples developed by different processing conditions following ASTM E-8 standard [152]. Uniaxial tensile tests were done at room temperature (25-30ºC) and cross-head velocity of 1 mm/min using Instron® 8862 servo-electric test system (10 ton). An extensometer of 25 mm gauge length was attached for the accurate measurement of strain. Three (3) tensile tests were conducted for each processing condition to ensure the repeatability in tensile properties. Tensile creep tests were carried out in constant load mode with initial stress level of 150 MPa at three different test temperatures: 550°C ± 2°C, 600°C ± 2°C and 650°C ± 2°C. Creep tests were done with the help of an Applied Test Systems (ATS), USA tensile creep testing machine having lever-arm ratio of 20:1. An extensometer with LVDT sensor was used for the accurate measurement of displacement or strain. At least two specimens were tested to assess the repeatability of creep tests. Specimen orientations with respect to rolled sample plate are shown in Fig.

3.6. The sample dimensions of uniaxial tensile and tensile creep specimens are shown in Fig. 3.7 and Fig 3.8, respectively.

Fig. 3.7: Schematic diagram showing the dimensions of sub-size tensile Specimen.

Fig. 3.8: Schematic diagram showing the dimensions of tensile creep Specimen.

3.7

Instrumented Charpy impact testing Full size Charpy V-notch specimens (10mm × 10mm × 55mm) were

fabricated from the as-received, heat treated and different thermo-mechanically processed samples at T-L orientation following ASTM E 23 standard [153], Fig. 3.6 and Fig. 3.9.

Fig. 3.9: Schematic diagram showing the dimensions of full size Charpy impact sample. 49

Chapter 3  

The Charpy impact tests were carried out over a range of temperature (-196 °C to +80 °C) to determine the ductile-to-brittle transition curves. The desired testing temperature was achieved by mixing different proportion of methanol and liquid nitrogen in solution bath. Specimens were soaked at desired temperature for at least 15 min before testing and the bath temperature was maintained within ±2 °C, using a thermo-couple for monitoring the temperature. The time span between removing the specimens from the bath and the impact test was kept within 5 s according to the ASTM E- 23 standard [153] for Charpy impact tests. All Charpy testing were conducted using an Instron 400 J impact machine (Model: SI-1C3), attached with an Instron® Dynatup Impulse Data acquisition system. The

Charpy impact

transition

curves

were

constructed

over

the

experimentally obtained scattered data points, following the procedure given by Sakai et al. 1983 [154]. The range of impact toughness and impact transition temperatures were measured from the impact transition curves obtained by joining the highest and lowest impact energy values as shown in Fig. 3.10.

Fig. 3.10: Schematic diagram showing the deviation in estimation of DBTT ( ± ΔDBTT) and USE (± ΔUS) values from tanh curve fitting.

3.8

Fractography Fractographic study has been conducted under SEM to evaluate the nature

of fracture, as well as, to locate cleavage initiation sites and void initiation sites on the fracture surfaces of broken Charpy specimens. The area close to notch / crack tip (upto 1 mm distance) was examined by tracing the cleavage river patterns back to the fracture origin. Energy Dispersive Spectroscopy (EDS) analysis has been employed to identify the nature of the inclusions. The microstructural unit over which cleavage crack propagates in an uninterrupted fashion is reflected by the cleavage facet size. Average facet size was determined based on the measurement of hundred (100) individual cleavage facets, from each sample tested at lower shelf region, in terms of the average of maximum and minimum dimensions.. 50

 

Results and discussions

 

 

Chapter 4 Effect of microalloy precipitates on microstructure and texture of hotdeformed modified 9Cr-1Mo steel

 

Chapter 4 4.1

Introduction and objective The modified 9Cr-1Mo steel emerged as a suitable material for in-core

applications in fast breeder reactors, owing to their high void swelling resistance under irradiation backed by high temperature strength, high thermal conductivity and low thermal expansion [155,156]. Modified 9Cr-1Mo steel is made superior to conventional 9Cr-1Mo steel in terms of creep properties by controlled addition of V and Nb, which form finely dispersed and stable MX precipitates (M stands for Nb and/or V, and X represents C and/or N) during tempering treatment of air-hardened martensitic structure. Fine MX precipitates preferentially form along the martensitic lath boundaries and stabilize the lath structure. Thus lath coarsening and the resultant structural softening upon prolonged thermal exposure is prevented, which improves the elevated temperature properties (especially creep and fatigue) of modified 9Cr1Mo steel [149,150,157,158]. Now, in order to be used as either steam pipes for ultrasupercritical (USC) power plants or fuel tubes in fast breeder nuclear reactors, heat resistant modified 9Cr-1Mo steel is processed through various hot-working techniques like forging, rolling and extrusion. Therefore, several studies have been carried out on the hot-deformation behaviour of 9Cr-1Mo steel [142–145,159–164]. Hot-deformation simulation experiments have been conducted using different techniques such as, uniaxial

compression,

plane-strain

compression

and

hot-torsion

[143,145,159,160,162–164]. The effect of processing parameters such as, strain (Ɛ), strain rate (Ɛ) and temperature (T) on the hot-flow behaviour, microstructural softening mechanism (recovery, recrystallization etc.), microstructural instability and defect generation have been studied [160,162,164]. Based on the results of such studies, deformation mechanism maps and constitutive models have been proposed for the representation of high temperature deformation of both conventional and modified 9Cr-1Mo steel [159,161,164]. Earlier studies on hot-deformation of modified 9Cr-1Mo steel, however, hardly investigated the following aspects: 

Effect of microalloy precipitates especially Nb precipitates, on the microstructural refinement and microstructural softening,



Effect of processing parameters and microalloy precipitates on the evolution of crystallographic texture. Present study is undertaken to address the above aspects.

51

Chapter 4  4.2

Results

4.2.1 Study of hot flow behaviour The hot flow curves of the Gleeble tested samples are presented in Fig. 4.1. In case of single-pass deformation schedules, decrease in deformation temperature increased the flow stress in almost a consistent fashion. At deformation temperatures of 1223K (950°C) and higher, the flow curves dropped slightly after reaching a peak value as indicated by arrows in Fig. 4.1a, which could be an indication of the occurrence of dynamic recrystallization, DRX, of prior austenite () grains [165–167]. Flow curve of 1173K (900°C) deformed sample does not show any drop, rather the flow stress increased initially with strain till a constant steady state level is reached. Such a flow curve is expected when dynamic recovery takes place without any DRX [165,166]. The variation of mean flow stress as the function of applied strain for the multi-pass (MP) deformed schedules are presented in Fig. 4.1b. In case of MP1 and MP2 schedules, flow stress increases initially with strain, i.e. with the decrease in deformation temperature, but decreases again at high strain level, Fig. 4.1b. MP3 schedule on the other side, showed an opposite trend as the flow stress continuously increased with the applied strain. In order to understand the effect of deformation schedules on the flow stress, microstructures of the hot-deformed samples need to be studied.

Fig. 4.1: (a) Flow curves for the single pass deformed samples and (b) the variation in mean flow stress with strain for the multi-pass deformed samples. 4.2.2 Microstructural study: Optical micrographs of single-pass and multi-pass deformed samples are shown in Fig. 4.2 and Fig. 4.3, respectively. The EBSD boundary maps showing the highangle boundaries (>15° misorientation indicated in black) and low-angle boundaries (215° misorientation indicated in red) inside single-pass and multi-pass deformed samples

52

Chapter 4 are given in Fig. 4.4 and Fig. 4.5, respectively. Comparing between the optical micrographs and EBSD maps and delineating the prior- grain boundaries, the average prior- grain size and grain aspect ratio for each sample are determined following the image analysis technique. Those are plotted in Fig. 4.6 as the function of finish deformation temperature (FDT).

Fig. 4.2: Optical micrographs of single pass (SP) deformed samples. Sample codes are mentioned on the images.

Fig. 4.3: Optical micrographs of multi pass (MP) deformed samples. Sample codes are mentioned on the images.

Fig. 4.4: EBSD micrographs showing high angle boundaries (black coloured) and low angle boundaries (red coloured) in single pass deformed samples. Sample codes are given on the images.

53

Chapter 4 

Fig. 4.5: EBSD micrographs showing high angle boundaries (black coloured) and low angle boundaries (red coloured) in multi pass deformed samples. Sample codes are given on the images.

Fig. 4.6: Variation in (a) average grain size and (b) grain aspect ratio as the function of finish deformation temperature in single pass and multi pass deformed samples. Abbreviations: PAG: Prior austenite grain size, Eff: Effective grain size, SP: Single pass, MP: Multi pass. Following major observations can be made from the microstructural study and EBSD analysis:



In case of single-pass (SP) deformed samples equiaxed prior- grain structures have been obtained in SP-1323, SP-1273, and SP-1223 samples, whilst in SP-1173 sample, the prior- grains were heavily elongated along the direction of metal flow i.e. perpendicular to direction of the applied strain, Fig. 4.2 and Fig. 4.4.



The  grains in SP-1323 sample were equiaxed but coarse. Grain size decreased significantly but the grains remained equiaxed as the FDT dropped to 1223K (950°C), Fig. 4.2, Fig. 4.4 and Fig. 4.6. Highly elongated grains with high aspect ratio were coarse again in SP-1173 sample, Fig. 4.2, Fig. 4.4 and Fig. 4.6. 54

Chapter 4 

The average density of high-angle boundaries, HAB, (length of boundaries per unit area) increased from ~2.8/ µm2 to ~3.1/ µm2 with the decrease in FDT from 1323K (1050°C) till 1223K (950°C), but the density of low-angle boundaries, LAB, remained quite similar (~3.0/ µm2), Fig. 4.4. The density of HAB and especially, LAB increased significantly in case of 1173K (900°C) deformed SP1173 sample to ~4.7/ µm2, Fig. 4.4.



The microstructures and boundary distributions are quite homogeneous in SP1323, SP-1273 and SP-1223 samples as those were comprised of recrystallized prior-γ grains. SP-1173 sample, on the other hand, indicated heterogeneous strain distribution in prior-γ structure. Intense flow localization within the deformation bands appearing in dark contrast inside the microstructure, indicated in Fig. 4.2d, and having high density of HABs and LABs, indicated in Fig. 4.4d. Formation of the deformation bands is a manifestation of flow instability, which can lead to failure and hence, is known to be detrimental for bending or forming point of view.



In case of multi-pass (MP) deformed samples: primarily equiaxed and recrystallized prior-γ grain structure with relatively small density of HAB and LAB (1.5 - 2.5/ µm2) have been found in all cases except MP1 schedule, Fig. 4.3 and Fig. 4.5.



Although some recrystallized grains were present, prior-γ grain structure of MP10.8 sample was primarily unrecrystallized and elongated in nature with high aspect ratio, Fig. 4.3, Fig. 4.5 and Fig. 4.6. Heavily deformed regions showed banded appearance, having high density of HABs and LABs (3.9 – 4.2/ µm2), as indicated in Fig. 4.5c. Such deformation bands are expected to form as a result of flow localization and grain subdivision. Among the multi-pass deformed samples, MP1-0.8 also showed coarse prior-γ grain size and the highest grain aspect ratio, Fig. 4.6.



The average prior-γ grain size of MP samples showed similar trend as followed by SP samples, i.e. the decrease in FDT from 1273K (1000°C) to 1223K (950°C) slightly decreased the grain size, but prior-γ grain size increased again at 1173K (900°C). MP1-0.4 and MP2-0.8 showed the most refined microstructures among the multi-pass deformed samples, Fig. 4.3(b,d) and Fig. 4.5(b,d).

55

Chapter 4  As per earlier studies on martensitic steels, prior-austenite grain boundaries and martensitic packet boundaries are high-angle boundaries, whereas, lath boundaries are the low-angle boundaries [25,27,32,168]. Based on the fact that the LABs are ineffective in retarding or deflecting the cleavage crack propagation and the HABs are only effective from that respect [28], the ‘effective grain size’ (EGS) can be determined from the EBSD analysis considering only the HABs (>15° misorientation). In martensitic microstructure, EGS primarily represents the martensitic packet size, which is an important microstructural parameter from mechanical property point of view. The EGS measured from the EBSD analysis for single-pass (SP) and multi-pass (MP) deformed samples, represented in Fig. 4.6a, showed finer EGS values in SP samples, apart from the FDT of 1173K (900°C). Hence, with respect to prior-γ grain sizes, the martensitic packet sizes appeared to be more refined in SP samples, in comparison to MP samples. That could be a beneficial effect of applying heavy singlepass deformation in comparison to multiple lighter deformation passes. The formation of deformed bands as a result of flow localization and its consequence on the distribution of crystallographic texture components can be analysed from the inverse pole figure maps with respect to the normal direction, i.e. the direction of applied strain (ND-IPF), given in Fig. 4.7.

Fig. 4.7: Typical EBSD inverse pole figure maps with respect to normal direction (NDIPF) of selected samples (sample codes mentioned) used for the determination of ODF. Colour legends for IPF maps are inserted. 4.2.3. Micro-texture study on hot-deformed samples As the deformed region in Gleeble tested sample is quite small, it is difficult to carry out macro-texture study (based on X-ray diffraction) over that region. Therefore, the distribution and intensity of different texture components in the deformed region was determined form the micro-texture study based on EBSD scan over large microstructural area from each sample (schematically shown in Fig. 4.8) so that 56

Chapter 4 statistically sufficient number of grains (required for the reliable determination of ODF) can be covered in the analysis. The inverse pole figure maps with respect to normal direction (i.e. ND-IPF maps) of the selected samples are shown in Fig. 4.7. It has been noticed that the different texture components were uniformly distributed in the recrystallized prior-γ grain structure of the single-pass deformed samples for higher FDT [≥ 1223K (950°C)], Fig. 4.7a. SP-1173 sample and also MP1-0.8 samples, on the other hand, showed the evidence of deformed prior-γ grain structure and inhomogeneous distribution of texture as elongated texture bands of //ND (in ‘blue’) and //ND (in ‘green’) were predominantly present almost in alternate fashion, Fig. 4.7(b, c). Such an arrangement of texture bands indicates the presence of deformed (un-recrystallized) condition of the prior-γ grain structure [169–172]. According to a recent study by Ghosh et al. [173], alternate arrangement of different texture bands can make steel prone to cracking (especially under impact loading) as strain incompatibility between different texture bands can initiate ductile crack along the interface of those bands. Multi-pass deformation following MP2 and MP3 schedules, although developed coarser structure compared to SP-1273 or SP-1223 samples, the texture components were uniformly distributed without the formation of any elongated texture band. That is certainly desirable from toughness and formability point of view.

Fig. 4.8: A typical Gleeble-deformed sample. RD, TD and ND stand for deformation direction (i.e. simulated rolling direction), transverse direction and normal direction, respectively. The φ2=45° ODF sections for selected single-pass and multi-pass deformed samples are presented in Fig. 4.9 and Fig. 4.10, respectively. Typical orientations found on the φ2=45° ODF section of Euler space for ferritic steel are shown in Fig. 4.9e as reference. From Fig. 4.9, it is evident that the texture of SP samples deformed at higher deformation temperatures [≥ 1223K (950°C)] comprised of predominantly rotated cube components ({001}), rotated Goss ({110}) and Goss ({110}) components. This is an indication that the prior-γ grains in those samples (just before martensitic transformation) were recrystallized grains [174], which turned into rotated 57

Chapter 4  cube and rotated Goss after martensitic transformation. The maximum texture intensity showed a slight decrease as the FDT dropped below 1273K (1000°C) which could be attributed to the lack of complete dynamic and metadynamic recrystallization of deformed γ- grain structures. Intensity of the above mentioned texture were rather weak in SP-1173 sample and a dominance of gamma-fibre (or ND-fibre) texture having //ND along with some other components like {332} and {554} could be noticed in this sample. Alfa fibre texture having //RD was also present in SP-1173 sample besides gamma fibre. The final texture in SP-1173 sample point to the fact that plane-strain deformation possibly developed copper ({112}) and brass ({110}) texture components in deformed γ-grains which directly transformed into martensite without recrystallization [174].

Fig. 4.9: (a-d) Orientation distribution function (ODF) at φ2=45° sections (Bunge notation) of single pass (SP) deformed samples. Samples codes are mentioned on the maps; (e) φ2=45° ODF sections showing the ideal orientations as observed in ferritic steels (BCC). Abbreviations: Rot-C: Rotated cube; Rot-Goss: Rotated Goss. Considering the effect of deformation temperature on texture, multi-pass deformed samples showed the exactly opposite trend as single-pass deformed samples, Fig. 4.10. Considering the effect of applied strain in MP1 schedule on the generation of texture, rotated cube texture was prominent in MP1-0.2 sample, Fig. 4.10a, which changed to a combination of alpha-fibre, gamma-fibre and small amount of rotated Goss texture in MP1-0.4 sample, Fig. 4.10b. Finally as the applied strain increased to 0.8, rotated cube and rotated Goss textures were completely replaced by RD-fibre (alpha) and ND-fibre (gamma) textures and other associated components, like {332}, Fig. 4.10c. Formation of such texture is expected as the deformed γ-grain

58

Chapter 4 structure transforms into martensite without recrystallization. This is an indication that small amount of recrystallization that happened after lighter deformation at the initial passes was completely suppressed after final deformation was applied at ~1273K (1000°C) and deformed γ-grain structure transformed into martensite. As the FDT decreased to 1223K (950°C), MP2-0.8 sample showed a combination of rotated cube texture and gamma-fibre texture, Fig. 4.10d. This suggests that the γ-grain structure was possibly partially recrystallized after final deformation pass as γ transformed into martensite. In contrary, after final deformation at 1173K (900°C), MP3 sample consisted of rotated cube and slight rotated Goss texture, (Fig. 4.10f), which are the characteristic feature of transformation from recrystallized austenite. Considering the effect of strain in MP3 schedule on its texture, initial passes developed a combination of rotated cube, alpha-fibre and weak gamma-fibre texture, Fig. 4.10e. It can be anticipated that the γ-grain structure was partially recrystallized (with retained strain) after initial deformation passes (at higher temperatures) and becomes completely recrystallized after applying the final pass at 1173K (900°C), this is when the accumulated strain is sufficiently large to drive the process of recrystallization to completion.

Fig. 4.10: Orientation distribution function (ODF) at φ2=45° sections (Bunge notation) of multi pass (MP) deformed samples. Samples codes are mentioned on the maps. 4.2.4. Precipitation study in hot-deformed samples TEM study on the hot-deformed samples along with EDS and SAED analysis identified the presence of different types of precipitates which are displayed in TEM images in Fig. 4.11 and are mentioned below in details. The precipitates in Fig. 4.11 are indicated by arrows. The size and volume fraction of the precipitates were 59

Chapter 4  determined from the TEM images by image analysis technique, assuming a constant foil thickness of ~ 50 nm. 

Ti-rich coarse precipitates (150-300 nm), containing N and small amount of Nb, were found at a low volume fraction of ~(5 – 10)×10-5 primarily at the prior-γ grain boundaries in all samples, Fig. 4.11a. SAED analysis indicated that the precipitates had a cubic crystal structure with lattice parameter of ~ 4.2 Å (i.e. close to ideal lattice parameter of TiN) and they followed ppt // α̍ type orientation relationship with the martensitic matrix, Fig. 4.11a. Such precipitates were possibly (Ti, Nb)N or (Ti, Nb)(C,N) precipitates, which were rich in Ti and N content.



The second group of coarse precipitates (180 – 380 nm) were present at relatively higher volume fraction of ~ (6 – 8) ×10-4, primarily along the martensitic lath and packet boundaries, Fig. 4.11b. Those precipitates were rich in Cr and had a polyhedral shape, Fig. 4.11b. Such precipitates were expected to be Cr23C6, which were commonly observed in 9Cr-1Mo steel.



Needle shaped Nb-rich fine precipitates (also containing V) of length varying between 80 – 200 nm and thickness less than ~ 50 nm were observed within the martensitic laths at volume fraction in the range of ~ (2 – 4) ×10-4, Fig. 4.11c. Those precipitates were more frequently observed in the samples deformed at lower temperatures [≤ 1223K (950°C)]. Such precipitates had a cubic crystal structure with lattice parameter of 4.4 Å and followed ppt // α̍ type orientation relationship with the martensitic matrix. This kind of needle shaped precipitates (possibly complex (Nb,V)(C, N) precipitates) could be the outcome of the formation of ‘V-rich wings’ on two sides of fine and spherical Nb(C, N) precipitates that develop during the hot-deformation of the steel [175].



Cluster of numerous fine spherical V-rich or Nb-rich (V, Nb)(C, N) precipitates have been found to form along the martensitic lath boundaries and also on the dislocation within the martensitic laths (volume fraction, (1 – 2) X 10-4), Fig. 4.11d. In general, the precipitates that were present along the boundaries were slightly coarser than the precipitates that formed within the laths. This could be attributed to the higher growth rate of the precipitate situated at the boundaries assisted by faster grain boundary diffusion as compared to lattice diffusion or

60

Chapter 4 even dislocation pipe diffusion. Spherical precipitates followed ppt // α type orientation relationship, besides random orientation. Overall, the microalloy precipitates observed in the investigated steel samples followed different orientation relationships with the martensitic matrix that were similar to, but do not exactly obey the classical orientation relationship expected in microalloyed steels such as, cube-on-cube, Kurdjumov-Sachs and Baker-Nutting relationship [176,177]. Such a phenomenon is expected when the precipitates form on the dislocations generated by hot-deformation in austenite due to strain-induced precipitation (SIP). The random arrangement of atoms at dislocation core does not allow the precipitates to follow ideal lattice correspondence with the matrix. Hotdeformation can also rotate the austenite crystals and alter the ideal orientation relationship for the precipitates formed even before deformation.

61

Chapter 4 

Fig. 4.11: Transmission electron micrographs showing precipitates (indicated by arrows) and their corresponding EDS and SAED analysis in the deformed (a) SP-1273, (b) SP-1173 (c) MP3-0.8 and (d) MP3-0.6 samples of modified 9Cr-1Mo steel.

4.3

Discussion

4.3.1. Effect of precipitates on microstructure and texture In order to explain the formation of different types of precipitates detected by TEM study on modified 9Cr-1Mo steel sample, thermodynamic prediction of fraction, stability and variation in internal composition of the precipitates have been performed using Thermo-Calc® software, Fig. 4.12. As predicted by the Thermo-Calc®, a small fraction of Ti-rich (Ti, Nb)N precipitates are expected to form in austenite at high temperature [~ 1423K (1150°C)], and the precipitate fraction remains almost unchanged during the subsequent cooling of the steel. Often complex (Ti, Nb)(C, N) precipitates can also develop in microalloyed steel, by the formation of NbC ‘Cap’ on top of pre-existing cuboidal ‘TiN’ precipitate ‘Core’ [177]. Small Ti-rich precipitate fraction can be attributed to the low Ti content as Ti was present as a trace element in the investigated steel. Although, even present as a small fraction, Ti-rich precipitates can pin down the prior-γ grain boundaries and restrict the γ-grain growth. Thermo-Calc® prediction in Fig. 4.12 indicates that besides Ti-rich precipitate, two other microalloy precipitates i.e. Nb- and V-rich precipitates can form in the modified 9Cr-1Mo steel at a much higher fraction than Ti-rich precipitates. Although Nb-rich carbo-nitride precipitate (also containing some V) starts to form at ~ 1403K (1130°C), its fraction increases sharply with the decrease in temperature and a significant extent of this Nb-rich precipitate is predicted to form over the hotdeformation temperature range [1173K (900°C) – 1323K (1050°C)]. Therefore, the strain-induced precipitation of Nb(C, N) could have taken place in fine and spherical form as indicated by the TEM study reported in Fig. 4.11. V-rich precipitates, also

62

Chapter 4 containing Nb (actually (V, Nb)N type precipitates), are predicted to form at relatively lower temperatures [≤ 1223K (950°C)] inside γ, Fig. 4.12. Cluster of such precipitates has also been observed in TEM study. According to Thermo-Calc®, Cr23C6 precipitates (also containing Mo) are expected to form during cooling at ~ 1153K (880°C) and its fraction can rise steeply to a high level, Fig. 4.12. Although the volume fraction of Tirich and V-rich precipitates obtained from TEM study were in the same order as predicted from Thermo-Calc®, the actual volume fraction of Cr23C6 precipitates were far less than the predicted level. This is because the hot-deformed samples were studied in as-quenched condition without tempering and Cr23C6 primarily form during the tempering treatment. In case of Nb- and V-rich precipitates, strain-induced precipitation is expected to happen during hot-deformation, which is believed to have increased the actual precipitate fraction. Mo-rich M6C type precipitate ((Mo, Nb)6C to be specific) is also predicted to form during tempering at ~ 873K (600°C), and those precipitates have not been found in the as-quenched samples.

Fig. 4.12: Volume fraction of different precipitates predicted using Thermo-Calc® thermodynamic software as the function of temperature (centre) and the compositional variation within the respective precipitates (mentioned on the images) with temperature. Since Nb-rich and V-rich precipitates, primarily form by strain-induced precipitation (SIP), those precipitates are expected to interact with the recrystallization of hot-deformed γ and thereby, control the microstructure and texture of the samples. Now, considering the Nb-rich and V-rich complex precipitates as a combination of Nb(C, N) and VN [as VC primarily forms at a temperature below ~ 973K (700°C)] their SIP kinetics and interaction with the recrystallization of deformed γ can be predicted from Dutta and Sellars model [178] and its generalized version recently proposed by Medina et al. [179,180]. It is known that the SIP of Nb(C, N) retard the austenite recrystallization and increases the recrystallization stop temperature (TNR) of the hot63

Chapter 4  deformed γ. For the prediction of TNR in Nb-microalloyed steels, among the various models available in literature, Dutta and Sellars model is simple and is most widely recognized [178]. SIP of VN can also retard the recrystallization of deformed γ but not to the extent of Nb(C, N) [180–183]. Based on extensive experimental studies, Medina et al. [179,180] recently proposed mathematical expressions for the prediction of SIP of VN. They have also proposed expression for the determination of activation energy of deformation (Qdef) [180], as the function of steel composition. That approach can be used for the prediction of recrystallization kinetics, i.e. recrystallization start and end times at different deformation temperatures, for any microalloyed steel and hence, is applied in the present study for elucidating the behaviour of the modified 9Cr-1Mo steel. The recrystallization-precipitation-time-temperature (RPTT) diagram for Nb(C, N) precipitation in the investigated steel (predicted following Dutta and Sellars model) for two different applied strains (0.2 and 0.8) are presented in Fig. 4.13a and b, respectively. The average  grain size developed after soaking at 1403K (1130°C) (~ 60 µm) is measured experimentally, and that is used as the starting grain size (D0) for the prediction of precipitation start time (t0.5P) from the following equation:. .

where,

3 10 Ɛexp

Ɛ

.

exp

exp

and

.

Eqn. 4.1

.

It is assumed that the entire C, N and Nb present in the steel were dissolved in γ as the SIP of Nb(C, N) started. The strain-rate applied in Gleeble® simulation Ɛ  1/

is used for the calculation.

Fig. 4.13: The recrystallization-precipitation-time-temperature (RPTT) diagram for (a and b) Nb(C, N) for applied strain of 0.2 and 0.8; (c) RPTT diagram for VN for 0.2 and 0.8 applied strain. RPTT diagrams are calculated following Dutta and Sellars [178] and Medina et al. [179,180]. Abbreviations: Ps: Precipitation start time, Pf: Precipitation finish time, Rs: Recrystallization start time and Rf: Recrystallization finish time. The following conclusions can be derived from the predicted RPTT diagrams for Nb(C, N), Fig. 4.13a and b:

64

Chapter 4  As the applied strain increased from 0.2 (as used in each pass of MP schedules) to 0.8 (strain per pass used in SP schedules), the TNR temperature (indicated in Fig. 4.13(a,b) by arrows) decreased from ~ 1343K (1070°C) to ~ 1233K (960°C). This could be the reason behind obtaining a deformed prior-γ grain structure in SP1173 sample as the corresponding deformation temperature [1173K (900°C)] was clearly below TNR. In multi-pass schedule, on the other hand, lighter deformation passes (Ɛ = 0.2) were applied for which TNR was much higher and as a result deformed prior-γ grain structure has been found even in MP1 schedule [FDT ~ 1273K (1000°C)].  However, the question arises why deformed prior-γ grain structure has not been found in MP2 and MP3 schedules as the FDT were even lower [1173K (900°C)1223K (950°C)] than MP1 schedule and certainly lower than the predicted TNR temperature. There the importance of inter-pass time can be identified. In MP1 schedule the inter-pass time is only ~ 10 s, over which SIP could possibly have effectively retarded the recrystallization of deformed γ. This is also reflected from Fig. 4.13a. However, as the inter-pass time increased to 30 – 50 s, the SIP completes within say 5 – 10 s and its retarding effect also diminishes at some point of the interpass time and recrystallization of γ can proceed over the remaining time interval. Although, the samples are cooled down at a constant rate (1K/s) over the inter-pass time, the situation is not expected to be much different from that indicated by the Dutta-Sellars model [178] as the cooling is not very rapid during the inter-pass interval. Therefore, with respect to the effect of deformation temperature on microstructure and texture, a completely opposite trend has been noticed between SP and MP schedules. Considering the effect of SIP of VN on the recrystallization of deformed-γ, Fig. 4.13c can be referred, where the VN precipitation start (PS) and γ-recrystallization start (RS) times are predicted for modified 9Cr-1Mo steel at different temperatures for the two applied strain levels (0.2 or 0.8) following the expression proposed by Medina et al. [179,180]. Since VN precipitation occurs at a lower temperature as compared to Nb(C, N) precipitation, its interaction with γ-recrystallization is also expected at lower deformation temperatures typically in the range of 1123K (850°C) – 1173K (900°C). This temperature range is at the lower side of FDT as applied in the present study. Its effect may also be more prominent in case of SP schedule with Ɛ = 0.8 (specifically SP-

65

Chapter 4  1173 sample) as the precipitation starts immediately after deformation, unlike MP schedule with Ɛ = 0.2, where an incubation period of more than ~ 10 s is required for the onset of SIP of VN, Fig. 4.13c. This could have contributed to the retention of heavily deformed prior-γ grain structure in SP-1173 sample (along with high flow stress), besides SIP of Nb(C, N). However, since SIP of VN follows the SIP of Nb(C, N), the formation and effect of VN precipitation can be influenced by Nb(C, N) precipitation. Considering the effect of deformation temperature on the microstructure of modified 9Cr-1Mo steel, high deformation temperature resulted in large prior-γ grain size in SP schedule. This could possibly be due to the complete recrystallization of γgrains along with sufficient grain growth during the short time interval, immediately after the deformation and before quenching the sample. In the MP schedule, the formation of coarse prior-γ grain size (and higher grain aspect ratio) can be attributed to the lack of recrystallization of the deformed γ-grain structure. Equiaxed prior-γ grain structure with smallest average γ grain size, as well as effective grain size, has been obtained in the samples deformed (or finish deformed) at 1223K (950°C) both in SP and MP schedules (i.e. SP-1223 and MP2-0.8) due to complete recrystallization and small γ-grain growth expected at that (lower) temperature. SIP of Nb(C, N) (also VN to some extent) could not have effectively retarded the recrystallization of deformed γ at that temperature, but could have restricted the γ-grain growth by pinning the grain boundaries as indicated by the presence of fine precipitates along the prior-γ grain boundaries in Fig. 4.11d. Precipitate particles can also contribute to the formation of deformation bands by interacting with the matrix dislocations in hot-deformed samples [166,184,185]. As the heavily deformed prior-γ grain structure transformed into martensite that not only sharpened the transformation texture and formed gamma-fibre along with {332} type components (which is expected [52,174]), but also intensified the HABs and LABs in the transformed martensitic structure, Fig. 4.4, 4.5, 4.9 and 4.10. The intensification of HABs, LABs and overall texture can be attributed to the following reasons: 

Formation of the deformation bands inside the heavily deformed regions can intensify the HABs in the microstructure.

66

Chapter 4 

As the deformed γ transformed into martensite, the dislocation present inside γ are inherited by martensite, which along with the dislocations generated to accommodate the transformation strain, increases the overall dislocation density in martensite [186]. When the MS temperature is relatively high, especially in low-C steel containing micro-alloying elements like Nb and V, which tie up the C further to form precipitates, autotempering of martensite can contribute to the recovery of dislocations and the intensification of the LABs.



The density of HAB and LAB may also depend on the variant selection during martensitic transformation, and the retained strain inside the prior-γ structure can promote the variant selection [11,187]. Finally, it is important to understand the effect of substitutional solute (alloying)

elements like Cr and Mo in the investigated steel on the SIP, γ-recrystallization and the resultant microstructure and texture after martensitic transformation. Based on the earlier studies the following points can be made in this regard: 

Substitutional solute elements like Cr and Mo, when dissolved in γ, are expected to retard the recrystallization of γ, by imposing a strong solute drag effect [165,167,188–190].



Presence of substitutional solutes can slow down the strain-induced Nb(C, N) precipitation to some extent [188].



Presence of Cr and Mo in γ-solution is expected to retard the transformation of austenite, which can significantly sharpen the transformation texture [52].



In MP schedule, when frequent high temperature deformation passes are applied, V and Mo primarily remain dissolved in γ. Solute drag from those elements combined with Zener drag from Nb(C, N) SIP could have imposed strong retardation to γ-recrystallization. As lower temperature passes are applied, more V and Mo may precipitate out (even on top of pre-existing Nb(C, N) precipitates), reducing the extent of solute drag and allowing γ-recrystallization to take place.

4.4

Summary In order to maintain a low rolling (or forging) load during hot-deformation,

modified 9Cr-1Mo steel is generally processed by multi-pass schedules maintaining high deformation temperatures and small inter-pass times. However, such a processing can result in the formation of deformed prior-austenite grain structure which may

67

Chapter 4  remain even after the first stage of normalization and tempering treatment as indicated by the recent studies from the present authors [18,19]. As such structure is detrimental from the toughness and formability point of view, double austenitization treatment has been recommended for restoring the equiaxed prior-γ grain structure, which is beneficial for those properties [19]. In order to effectively design the thermomechanical processing schedule for developing fine and equiaxed prior- grain structure, present study reports important findings as listed below: 

The pinning effect from strain-induced Nb(C,N) precipitation along with the high solute drag effect from Cr, Mo and V, effectively retarded the recrystallization of deformed-γ at high deformation temperature and short inter-pass times (~ 10 s) which resulted in the formation of deformed γ-grain structure.



Application of multiple deformation passes maintaining sufficient inter-pass interval (30 – 50 s) can be beneficial for the formation of fine and equiaxed γgrain structure by static recrystallization.



As longer inter-pass interval is used, sufficient time remains for the static recrystallization of γ to take place after the completion of strain-induced Nb(C, N) precipitation. Such a deformation schedule however, can increase the rolling load at lower deformation temperatures (due to higher flow stress) and it can also develop cube texture in transformed martensite, which is considered to be detrimental for the steel from delamination and low-temperature toughness point of view [52]. If high capacity rolling mill is available for applying heavy deformation passes,

in that case high temperature deformation (preferably in single-pass) can be beneficial as that can be favourable for promoting dynamic recrystallization of deformed γ. Straininduced Nb(C, N) and also VN precipitates can retard the recrystallization of γ when single deformation pass is used at a lower temperature [~ 1173K (900°C)].

68

 

Chapter 5 Effect of hot-rolling temperature on the ductile-brittle transition behaviour of modified 9Cr-1Mo steel

 

Chapter 5 5.1

Introduction and objective The modified 9Cr-1Mo steel is generally used in hot-rolled, normalized and

tempered condition. The beneficial effect of normalizing and tempering treatment on the mechanical properties of this steel is widely investigated [15,19,27,32,191,192]. In general, normalization treatment by austenitizing the steel can have significant effect on both strength and toughness [193]. Normalization treatment helps in relieving the rolling strain, homogenizing the microstructure and dissolving the coarse carbide particles and thus improves the mechanical properties [194]. Application of repeated cycles of normalization treatment can be further beneficial as it can significantly refine the prior austenite grain size [194]. The beneficial effect of grain refinement in 9Cr-1Mo steel is well established and the refinement in prior-austenite grain size and precipitate size have been found to improve the strength and impact toughness of the steel over a wide range of temperature [15,19,27,29,32,191–195]. However, a continuous drive has always been there to improve the strength further, which can hamper the impact toughness. The microstructural refinement depends on the processing parameters and the rolling schedule is known to play an important role on the microstructure, precipitation and mechanical behaviour of steel [15,19,24,27,32,149,191–193,196–199]. As an example, Klueh et al. [149] achieved an improvement in high temperature mechanical properties by applying thermo-mechanically controlled processing of 9-12 wt% Cr containing steels. Compared to the studies carried out to understand the effect of heat treatment parameters, the effect of processing parameters during hot-working on the microstructure and properties of 9Cr-1Mo steel is far less investigated. Selected studies on this area also include the effect of austenitization treatment as the steel is normalized after hot-working [144,197]. Thus the sole effect of hot-working parameters on the mechanical properties, especially the strength and impact toughness of 9Cr-1Mo steel has hardly been studied and that sets the objective of the present work. Deformation temperature is known to influence the strength-impact toughness combination of lowcarbon steels having ferrite-pearlite microstructure [15,29,191,193]. This chapter aims to study the influence of deformation temperature on microstructural parameters, crystallographic texture and mechanical properties (tensile properties and impact transition behaviour) of modified 9Cr-1Mo grade martensitic steel.

69

Chapter 5  5.2

Results and Discussion

5.2.1

Characterization of microstructure Optical micrographs and SEM micrographs of the As-received and hot-rolled

samples showed predominantly lath martensite structure, Fig. 5.1, Fig. 5.2 and Fig. 5.3. There was no ferrite in the As-received steel as it was normalized at 1050C, which is much above the Ae3 temperature (889C, predicted from Thermo-Calc® software) of the steel. Deformation of austenite just above, around or below the TNR (~950C) possibly have reduced the hardenability of the steel, restricted the complete martensite transformation during air-cooling and resulted in the formation of ferrite, Fig. 5.3 and

Table 5.1. The amount of ferrite increased with the decrease in rolling temperature, Table 5.1. Inter-critical deformation inside austenite + ferrite region led to maximum amount of ferrite formation in 875-HR sample, Fig. 5.3 and Table 5.1. Decrease in rolling temperature has been found to reduce the prior-austenite grain size (PAGS) and increase the grain aspect ratio, Table 5.1. In case of high rolling temperatures, prioraustenite grains become more equiaxed and coarse due to recrystallization and grain growth. Decrease in rolling temperature restricted the grain growth, although, the presence of deformed and elongated grains in the partially recrystallized microstructures increased the grain aspect ratio, Fig. 5.2, Fig. 5.3 and Table 5.1. The martensitic packet sizes were in the close range in the investigated samples, Table 5.1. Normalization treatment contributed to the formation of relative small PAGS and martensitic packet size in the As-received steel, Table 5.1.

Fig. 5.1: (a) Optical and (b) SEM micrographs of the As-received steel.

70

Chapter 5

Fig. 5.2: Optical micrographs of the hot-rolled samples: (a) 1050-HR, (b) 1000-HR, (c) 950-HR and (d) 875-HR.

Fig. 5.3: SEM micrographs of hot-rolled samples: (a, b) 1050-HR, (c, d) 1000-HR, (e, f) 950-HR and (g, h) 875-HR; (a, c, e, g) before tempering and (b, d, f, h) after tempering treatment. Ferrite regions are indicated by arrows.

71

Chapter 5  Table 5.1: Microstructural parameters, fraction of low-angle boundaries and the effective grain sizes (estimated from EBSD analysis) and cleavage facet sizes (measured from fractographic study) of the As-received and hot-rolled samples. Sample Ferrite fraction (%) PAGS (µm) Aspect ratio

Asreceived -

1000-HR

950-HR

875-HR

1.5

3.1

6.8

12.3

43.7±24.1 67.4 ± 10.2 65.7 ± 13.7 49.3 ± 11.3 4.9

Martensite Packet size 6.6 ± 2.4 (µm) Volume fraction (%) of coarse M23C6 type 4.76 precipitates (>0.2 µm) Volume fraction (%) of fine MX type 0.042 precipitates (>50 nm) Fraction of low-angle 56% boundaries ( 2 µm) possibly nucleated from the Cr23C6 particles, was found to be lower in 1000-HR sample, than that in 950-HR sample, but

82

Chapter 5 these voids were larger in size in the former sample, Fig 5.14(c, d). Some large but uniform sized voids were present in As-received sample along with small void-sheet joining ligament. Small voids could have nucleated around the fine microalloy precipitates. These observations are in accord with the earlier discussion on the effect of precipitate fraction and size on upper shelf energy (USE).

Fig. 5.14: SEM fractographs of (a, f) As-received, (b, g) 1050-HR, (c, h) 1000-HR, (d, i) 950-HR and (e, j) 875-HR samples tested at the (a-e) upper shelf regime (test temperature of +80C) and (f-j) lower shelf regime (test temperature of -100C). Some of small voids are indicated by black arrows, while large voids are indicated by red arrows. Cleavage facets were measured on the fracture surfaces of the specimens tested at lower shelf regions and their average sizes are listed in Table 5.1. The average facet size was highest in 875-HR, followed by 1000-HR and 950-HR sample, Table 5.1. Asreceived steel and 1050-HR samples showed smaller facet sizes. Cleavage facet size

83

Chapter 5  provides an indication of the resistance faced by the cleavage crack during its propagation and smaller facet size represents higher resistance to cleavage fracture (as in As-received and 1050-HR samples). Higher undulation on the fracture surface and grater angular difference between the adjacent facets further explains the higher resistance of 950-HR sample to cleavage fracture than that of 1000-HR sample, Fig. 5.14. As per the results in Table 5.1 the average cleavage facet size (16-20 µm) was much smaller than the prior-austenite grain size (44-67 µm) and larger than the martensitic lath size (~0.1-0.6 µm) of the investigated samples. Hence, prior austenite grain size and lath size can not be the microstructural parameters that control cleavage fracture in martensitic steel. The similarity in values between cleavage facet size (1620 µm) and effective grain sizes measured from misorientation angle (angle-axis pair) between martensitic boundaries (~18-20 µm) signifies the importance of ‘effective grain size’ concept in explaining the cleavage fracture behaviour of martensitic steels. The difference (2-3 orders of magnitude) between the cleavage facet size and the martensite packet size (6.5-8.5 µm) might have appeared from the stereological factor involved in comparing a 3-Diemsional feature (facet size) and 2-Dimensional feature (packet size) or martensitic packets might not be the ‘effective grain’ in martensitic steel for improving cleavage fracture resistance. Any other factor related to the crystallography of martensite structure could also have contributed to such a difference and detailed EBSD analysis is needed to resolve this issue.

5.3

Summary Modified 9Cr-1Mo steel plate, received in normalized and tempered starting

condition, were subjected to hot-rolling treatment at four different temperatures; among which two were above TNR : 1050C (1050-HR) and 1000C (1000-HR), one around TNR: 950C (950-HR) and one was below TNR at inter-critical region: 875C (875-HR). Hot-rolled samples were finally tempered and tested for tensile and Chapy impact properties. 

Hot-rolling significantly increased the strength of the 9Cr-1Mo steel (by 150-200 MPa), sacrificing some amount of ductility (by 10-15%).



High rolling temperature (1050C) is preferable in order to achieve good combination of strength and ductility (YS800 MPa with 15% elongation). 84

Chapter 5 Combined effect of solid-solution strengthening and precipitation strengthening contribute to the high strength of 1050-HR sample. 

In case of intermediate rolling temperature (950C), precipitation strengthening and dislocation strengthening offered maximum strength (YS812 MPa), whilst, ductility reduced to minimum (12%).



Formation of high fraction of ferrite during inter-critical deformation and coarsening of precipitates during the final tempering treatment reduced the strength of 875-HR sample (YS686 MPa).



The best impact properties of the As-received steel could be attributed to the normalization treatment, which resulted in (i) small effective grain size, (ii) lowfraction of {001} cleavage planes parallel to the main fracture plane of the sample and (iii) high fraction of slip planes of the crystals parallel to the shear plane of the sample.



Hot-rolling increased the ductile-brittle transition temperature by 26-40C and also reduced the upper shelf energy (USE) of the steel.



Among the rolled samples, 1050-HR showed the best combination of DBTT and USE, which could have resulted from (i) relatively smaller effective grain size, (ii) beneficial texture such as, -fiber (//ND) and {554} component and (iii) low-fraction of harmful coarse precipitates.



Small effective grain size and beneficial texture could have offered high USE to 950-HR sample, however, presence of dislocation sub-structure and high matrix strength could have restricted the plastic deformation and increased the DBTT of that sample.



Formation of dislocation substructure and high ferrite fraction hampered both DBTT and USE of inter-critically rolled 875-HR sample.



Finally, if high strength is required in modified 9Cr-1Mo steel for certain application then the steel can be used in as-hot-rolled condition and it is preferable to roll at high temperature as it increases strength, with minimum negative effect on the ductility and impact toughness.

85

 

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86

 

Chapter 6 Effect of normalization temperatures on ductile-brittle transition behaviour of modified 9Cr-1Mo steel

 

Chapter 6 6.1

Introduction and objective In general the beneficial effect of grain refinement on the decrease in DBTT and

the

increase

in

USE

of

ferrite

/

martensite

steel

is

widely

reported

[15,27,32,78,191,192]. An inverse square-root dependence of DBTT with ferrite grain size or prior-austenite grain size has been suggested in many alloys, starting from pure iron and very low alloyed steel [15,191,192], microalloyed steel [15,78], and finally low-carbon martensitic steel [27,78]. Smaller prior-austenite grain size has also been reported to show better dynamic fracture resistance and lower DBTT in ferritemartensite steels [218,219]. More refined structure of tempered martensite was found to decrease the DBTT even in neutron irradiated condition, compared to coarser martensitic structure [220]. A fine grained tempered martensite structure is also desirable for obtaining good creep properties [131,221]. Though earlier literature indicated mostly the effect of ‘metallographic grain size’ on the ductile-brittle transition behaviour of ferrite / martensite steels [218,219], recent studies on the cleavage fracture are signifying the effect of coherent crystallographic microstructural unit, i.e. ‘effective grain size’ on DBTT of low-carbon ferritic steels [17,24–26]. The concept of ‘effective grain size’ on the fracture behaviour of martensitic steel is yet to be established. According to the recent studies on low-carbon ferritic steels, grain size is not the sole factor that determines the ductile-brittle transition behaviour under impact loading condition. Several other factors such as, distribution of grain boundary misorientation angles, precipitates, strength of the steel matrix and crystallographic texture can also influence the DBTT and USE [25,200,222]. The effect of all those factors on the mechanical properties has hardly been considered earlier in martensitic steel. Present study aims to investigate the combined effect of different factors that influence the impact toughness and impact transition behaviour of 9Cr-1Mo tempered martensitic steel with ultimate aim of minimizing the DBTT in un-irradiated condition. As the 9Cr-1Mo steel finally used in normalized and tempering condition in actual service condition, the beneficial effect of cyclic heat-treatment, especially the double austenitization treatment, on the austenite grain refinement and the resultant improvement in Charpy impact toughness of Cr-Mo type of steels has been studied in the earlier studies [27,168,223,224]. This chapter throws a fresh look on this aspect in 9Cr-1Mo steel and emphasised on the effect of austenitization temperature on the strength, ductility and impact toughness of this steel. 87

Chapter 6  6.2.

Results and Discussion

6.2.1

Characterization of microstructure Optical micrographs, Fig. 6.1, and SEM micrographs, Fig. 6.2, of the As-

received and reheated (i.e. double austenitized) samples showed the lath martensite structure throughout the matrix. No evidence of ferrite was found in any of the sample. Formation of complete martensite structure is expected in the investigated samples as the reheating temperatures (950-1100C) were higher than the Ae3 temperature (complete austenite formation temperature under equilibrium condition) of 889C as predicted from the Thermo-Calc® software. Effect of hot-rolling remained in the Asreceived sample even after the first stage of austenitization treatment as evident from the prior austenite grain structure, which was elongated along the rolling direction, Fig.

6.1(a) and Fig. 6.2(a, b). The prior austenite grains become more equiaxed after the double austenitization treatment, Fig. 6.1(b-d) and Fig. 6.2(c-h), which is reflected from the decrease in average grain aspect ratio, Table 6.1. The prior austenite grain size (PAGS) and martensite packet size also decreased after the double-austenitization treatment, Table 6.1. Among the reheated samples, 1100-Reheat sample showed the largest PAGS and packet size, followed by 1025-Reheat and 950-Reheat samples,

Table 6.1. The PAGS and martensitic packet size of 950-Reheat and 1025-Reheat samples were close to each other and within the standard deviation of the measured values, Table 6.1.

Fig. 6.1: Optical micrographs of As-received and reheated samples of 9Cr-1Mo steels: (a) As-received, (b) 950-Reheat, (c) 1025-Reheat and (d) 1100-Reheat.

88

Chapter 6

Fig. 6.2: SEM micrographs of (a, b) As-received steel and (c-h) reheated samples. Micrographs of the reheated samples are taken (c-e) before tempering and (f-h) after tempering treatment. Table 6.1: Microstructural parameters of the as-received and reheated samples. Fraction of low-angle boundaries and the effective grain sizes were measured from the EBSD analysis. Cleavage facet sizes were measured from the fractographic study. Sample

As-received

950-Reheat

1025-Reheat

1100-Reheat

Aspect ratio

4.9

1.1

1.4

1.3

PAGS (µm)

43.7 ± 24.1

12.1 ± 4.0

14.0 ± 4.4

18.6 ± 7.9

6.6 ± 2.4

3.4 ± 1.5

3.7 ± 1.5

4.6 ± 1.7

18.6

7.8

8.14

11.4

4.76

3.45

2.47

2.78

56%

75%

49%

42%

16.3 ± 4.1

6.7 ± 1.9

6.5 ± 1.8

9.6 ± 2.5

Martensite Packet size (µm) Effective grain size (µm) Precipitate volume fraction (%) Fraction of low-angle boundary ( 2 µm) nucleated from the Cr23C6 particles was found to be higher in Asreceived and 950-Reheat samples than the rest. Large fraction of small and uniform sized voids was found in 1025-Reheat and 1100-Reheat samples and those voids could have nucleated around the fine microalloy precipitates. These observations are in agreement with the earlier discussion on the effect of precipitate fraction and size on impact energy absorption under ductile region.

Fig. 6.12: SEM fractographs of (a, e) As-received sample, (b, f) 950-Reheat sample, (c, g) 1025-Reheat sample and (d, h) 1100-Reheat sample tested at the (a-d) upper shelf regime (test temperature of +40C) and (e-h) lower shelf regime (test temperature of 130C); (i) the energy dispersive spectroscopy (EDS) analysis of the particle indicated by arrow in (e). Average size of the cleavage facets measured on the fracture surface of the specimens tested at lower shelf regions are listed in Table 6.1. The average facet size was highest in As-received steel, followed by 1100-Reheat sample. 1025-Reheat and 950-Reheat samples showed smaller facet sizes. Cleavage facet size gives an indication of the resistance faced by the cleavage crack during its propagation and finer facet size indicates higher resistance against cleavage fracture (as in 1025-Reheat and 950-Reheat samples). Moreover, the undulation on the fracture surface and the angular difference

102

Chapter 6 between the adjacent facets are also higher in 1025-Reheat sample than the other samples, which further explains the higher resistance against cleavage fracture and lower DBTT of that sample. Similar to the previous study on hot-rolled specimens, the measured values of average cleavage facet size (6-16 µm) was much smaller than the prior austenite grain size (18-44 µm) and much larger than the martensitic lath size (0.1-0.4 µm) of the normalized samples. 2-3 order of magnitude difference between the martensite packet size (3-7 µm) and the cleavage facet size indicates the possibility that the packet size might not be the true representation of the ‘effective grain size’ for cleavage fracture resistance of lath martensitic steel.

6.3

Summary The As-received samples were subjected to reheating treatment (i.e. double

austenitization treatment) at three different temperatures namely, 950C (950-Reheat), 1025C (1025-Reheat) and 1100C (1100-Reheat) and the samples were finally tempered. The major findings derived from the detailed microstructural and precipitation study, EBSD analysis, macro-texture study and mechanical property evaluation (hardness, tensile and Charpy testing) of the samples are listed below: 

The double austenitization treatment converted the coarse and elongated prior austenite grain structure of As-received steel to finer and more equiaxed grain structure in the reheated samples. 1025-Reheat and 950-Reheat samples showed finer prior austenite grain size and martensitic packet size than the remaining samples (As-received and 1100-Reheat sample). This could suggest that the first normalization treatment (i.e. 1050ºC for 25 min) might be insufficient to completely eliminate the initial rolled microstructure and to dissolve the precipitates. If the first heat treatment was optimized in this regard, the necessity of the double austenitization treatment could be different.



EBSD analysis showed the highest fraction of low-angle boundaries (15°), the B5 fails to cause any substantial deviation in cleavage crack path. A few more examples given in Table 8.3 further confirm the greater effectiveness of block boundary over packet boundary in deviating the cleavage crack propagation. These results also highlight the necessity for consideration of the projection angle to get an insight of the cleavage crack propagation phenomenon. Moreover, the ineffectiveness of packet boundary was found, when EBSD study was carried out over the cleavage facets by manually collecting data points. The typical example shown in Fig. 8.9, indicates the existence of two variants V5 (A) and V15 (B) within a single cleavage facet, and that of two other variants V9 (C) and V14 (D) in an adjacent facet. Therefore, the boundaries between these two sets of variants are the packet boundaries, which were unable to resist the cleavage crack propagation within a given facet on the main fracture plane. The active cleavage planes are given in the parentheses in Fig. 8.9. Corresponding misorientation, cleavage and projection angles are listed in Table 8.3.

140

Chapter 8

Fig. 8.9: EBSD study of cleavage facets in fracture surface. Table 8.3: Few examples of interacting martensitic boundaries and cleavage crack given in cracked microstructures and cleavage facet. Site number

Intervariant boundary

Misorientation angle (°)

Cleavage angle (°)

Active cleavage planes

Angle of deviation (measured) (°)

Cleavage Projection angle (°)

Type of boundary

Site

136-137

V3-V19

15.4

5.2

{010}-{010}

3

1.2

Packet

Fig.

137-138

V19-V23

60

39.8

{010}-{001}

40

37

Block

8.7c

152-153

V17-V18

58.9

46.6

{100}-{001}

2

0.2

Block

153-154

V18-V17

59.2

47.3

{001}-{100}

1

1.3

Block

154-155

V17-V24

52.1

31.3

{100}-{100}

24

22

Packet

164-165

V11-V15

57.6

48.7

{010}-{001}

27

155.2

Packet

164-166

V11-V10

58.0

40.8

{100}-{100}

38

143.6

Block

72-73

V8-V7

56.8

39

{100}-{100}

38

35.4

Block

Fig.

73-75

V7-V12

48.2

33.1

{100}-{010}

2

0.7

Block

8.7b

75-78

V12-V8

58.9

41.0

{010}-{001}

40

38.2

Block

83-84

V8-V7

57.2

38.6

{010}-{001}

43

38.4

Block

84-85

V7-V10

10.4

6.7

{001}-{100}

0

0.7

Sub-block

85-86

V10-V7

9.3

4.1

{100}-{010}

4

2

Sub-block

A-B

V5-V15

20.8

5.2

{100}-{100}

-

-

Packet

Fig.

C-D

V9-V14

14.1

2.2

{100}-{010}

-

-

Packet

8.9

In order to further explain the interaction between the cleavage crack and the boundaries situated inside the martensitic structure, a schematic diagram is shown in

Fig. 8.10. In this case, the actual orientations of three arbitrary crystallographic variants (V3, V9 and V10) are chosen with respect to the sample orientation. In spite of being a

141

Chapter 8  packet boundary, the interface between V3 and V9 variants is expected to cause a small deviation (6.07°) in the cleavage crack. In contrast, the block boundary between V9 and V10 variants can enforce effective deflection of the crack path by 48.19°.

Fig. 8.10: Schematic representation of cleavage crack propagation across three different martensitic variants. 8.3.4

Hierarchical martensitic microstructures and high temperature dynamic

deformation In order to identify the ability of different sub-block, block and packet boundaries to provide ductile fracture resistance, the angular deviation between slip planes of adjacent martensitic variants has been calculated. Interestingly, the result shows that there can be very few combinations of inter-variant boundaries (i.e. some packet boundaries) that follow high angular deviation (> 15°). For example, if {110} planes are considered as the active slip planes, high-angular deviation is expected only in case of ~18.5% of all inter-variant boundaries, Fig. 8.11a, whilst that possibility disappears completely if both {110} and {112} types of planes are considered together as the slip planes, Fig. 8.11b and Table A3 (in Appendix section). Hence, there is no clear indication about any hierarchical martensitic boundaries (e.g. packet, block and sub-block) that can be considered as effective as block boundaries in terms of cleavage crack deflection. Moreover, angular deviation between the slip planes are not the only determining factor for slip to occur, rather the angle between loading direction and respective slip direction ( type) (that decides the Schmid factor) also plays an important role. Therefore, calculating Schmid factors of the different martensitic variants, without considering the local arrangements of crystallographic planes, cannot provide proper understanding on dynamic ductile cracking phenomenon of martensitic steel.

142

Chapter 8

Fig. 8.11: Distribution of angles between slip planes for all combination of martensitic variants considering (a) {110} as active slip system, and (b) both {110} and {112} as active slip systems. The microstructural features on the cross-sectional surface (RD-TD plane) of the high-temperature impact fractured samples indicated the presence of adiabatic shear bands (ASB), Fig. 8.12. Primarily two types of ASBs are observed: (a) straight bands (indicated by yellow arrows) having width of 100-200 µm, and (b) rotated bands (shown by red arrows) having width of 30-50 µm. Shear cracks are found to have formed more frequently at the interfaces of rotated ASBs (17.3 ± 3.7 per 1000 µm2 area), as pointed by the black arrows in Fig. 8.12, as compared to the straight ABSs (4.1 ± 2.4 per 1000 µm2 area). This difference in shear cracking phenomenon within ASBs as well as the dimension of ASBs (30-150 µm) primarily indicates that the size and distribution of prior-austenite grains could have a significant influence on this phenomenon as compared to the other hierarchical martensitic structural units.

Fig. 8.12: Optical micrographs of adiabatic shear band (ASB) zones in 1000HR1100NT samples fractures at high temperature (+75°C). Rotated and straight ASBs are indicated by indicated by red and yellow arrows, respectively, whilst cracks at rotated ASBs are shown by black arrows. Therefore, four different samples (Fig. 8.13) with the following combinations were chosen for detailed study on the deformation mechanism responsible for the formation of ASBs and the related cracking phenomenon:

143

Chapter 8  

1000HR-950NT: Very small PAGS (13.2 ± 4.1 µm) and very narrow grain size

spread (10-25 µm), Fig. 8.13a. Showing high USE (223J). ‘Sample code: S1’.



875HR-1100NT: Small PAGS (27.1 ± 4.9 µm) and narrow grain size spread (12-

40 µm), but both are almost two times larger than S1, Fig. 8.13b. Showing low USE (74J). ‘Sample code: S2’.



1050HR-1100NT: Large PAGS (61.6 ± 23.7 µm) and wide grain size spread (20-

125 µm), Fig. 8.13c. Showing high USE (237J). Sample code: L1.



1000HR-1100NT: Large PAGS (60.3 ± 39.1 µm) almost similar to L1, but very

wide grain size spread (25-275 µm) that is nearly two times larger than L1, Fig. 8.13d. Showing high USE (191J). Sample code: L2. Careful microstructural examination and prior-austenite grain size distribution reveal the presence of ‘mixed grain structure’ comprising of both large and small grains in L1 and L2 samples. The fraction of areas covered by the prior-austenite grains having sizes larger than the mean PAGS are found to be 67% in L1 and 85% in L2, Fig. 8.13(c and d).

Fig. 8.13: Optical micrographs showing prior-austenite grain (PAG) structures of (a) 1000HR-950NT (S1), (b) 875HR-1100NT (S2), (c) 1050HR-1100NT (L1) and (d) 1000HR-1100NT (L2) samples. (e) PAG size distributions of these samples. 8.3.5

Evolution of crystallographic texture within adiabatic shear bands EBSD-IPF maps in Fig. 8.14 show that shear cracking occurred within the

ASBs constituted of alternate bands of gamma-fibre ({111}//ND) and Goss-fibre ({110}//ND) textures, Fig. 8.14. This observation is similar for all the four samples selected for this investigation. However, some cube texture bands (i.e. {001}//ND) that are known to be detrimental from ductility and toughness point of view were 144

Chapter 8 occasionally found at different distances (~ 20-50 µm) from the shear cracks, Fig.

8.14(a-c). Presence of cube texture band at the close vicinity of shear cracks in S2 sample could have restricted its plastic deformation during dynamic loading, resulting in low USE (74 J).

Fig. 8.14: IPF maps showing adiabatic shear bands (ASBs) adjacent to ductile shear cracks of (a) S1, (b) S2, (c) L1 and (d) L2 samples (a-d counter clockwise sequence). IPF legend is given in inset. In order to understand the underlying deformation behaviour, the evolution of the texture in and around the shear band zones of the aforementioned four samples was simulated and compared with the experimental results. Simulation of texture evolution was performed by imposing a velocity gradient tensor for simple shear deformation using the Visco-plastic self-consistent polycrystal model, as explained in Section 8.2.2. The experimental texture was represented by 10000 discretised crystal orientations and simple shear deformation with a shear strain-rate  was applied to it as defined by the following velocity gradient tensor in a rectangular reference system:

 0  0    L   0 0 0  0 0 0  

Eqn.17

The strain-rate increment step used for this analysis was 0.005. The amount of shear strains (  / 2 ) considered in the simulations were 0.5, 0.38, 2.05 and 2.2 for the S1, S2, L1 and L2 samples, respectively. The strain values were obtained

145

Chapter 8  experimentally by averaging the slopes of the adiabatic shear bands with respect to the specimen axis (i.e. RD in the present investigation) following Moss’s method [315]. After comparing the simulated texture to the experiment it was found that best texture predictions were obtained employing different interaction parameters in the VPSC polycrystal code: (i) ‘Secant model’ for S1 (grain size: 13.2 ± 4.1 µm), (ii) ‘Tangent model’ for S2 (grain size: 27.1 ± 4.9 µm) and L1 (grain size: 61.6 ± 23.7 µm) and (iii) relative directional compliance (RDC) model [316] for L2 (grain size: 60.3 ± 39.1 µm). Obtaining best simulated texture results with the help of different grain interaction parameters for different samples of the same material is quite unusual. However, it has been found already that the interaction parameter can be dependent on grain shape [317], and grain size [318–320]. The initial microstructures and textures (before impact loading) of the four selected samples were different owing to different rolling temperatures (with respect to TNR) and different normalizing temperatures, Fig. 8.15(a1, b1, c1 and d1). However, after dynamic deformation, the pole figures of all the specimens primarily show typical bcc shear textures, i.e. {110} // shear plane and // shear direction with respect to the loading geometry. The intensity of {110} and {111} poles of these samples are strong along the TD (i.e. shear plane normal, SPN) and RD (i.e. shear direction, SD), respectively, see Fig. 8.15(a2-a3), (b2-b3), (c2-c3) and (d2-d3). Similar to the pole figures, the orientation distribution functions (ODFs) of the deformed samples primarily indicate the intense presence of ideal shear texture components:

{111}  121  and {111}  112  as shown in the respective ODFs, Fig. 8.15(a2-a3), (b2b3), (c2-c3) and (d2-d3). Apart from these major texture components, the presence of {1 1 0}  0 0 1 

and

{1 1 0 }  1 1 1 

components in L1, and

{1 1 0 }  1 1 1 

components in

L2 have also been observed, Fig. 8.15(c and d). Presence of these texture components are in agreement with the earlier report on the dynamic deformation of pure iron [282]. The presence of

{1 1 0 }  1 1 1 

components in deformed samples indicate the

occurrence of dynamic recrystallization within adiabatic shear band regions [317–320].

146

Chapter 8

147

Chapter 8 

Fig. 8.15: 1: Initial, 2: experimentally obtained, and 3: simulated textures (i.e. pole figures of (110), (111) and (112), and ODF at φ2=45°) of (a) S1, (b) S2, (c) L1 and (d) L2 samples. 8.3.6

Rotational effects within adiabatic shear bands

 

Crystal plasticity calculations suggest that dynamic loading leads to a marginal

increase in the rotation of ellipsoids (i.e. deforming grains) within the aggregate (i.e. adiabatic shear band in this case) of S1 (up to ~0.5°) and S2 (up to ~1.5°) specimens,

Fig. 8.16. The extent of average grain rotations were higher in L1 (up to ~41°) and L2 (up to ~28°) specimens, Fig. 8.16. Similar phenomenon of lower amount of rotation in smaller grains, and higher rotation in coarse grains are also observed in dynamically deformed high purity Cu [294].

Fig. 8.16: Variation in average grain rotation within adiabatic shear band regions of the investigated samples. The grain rotation is defined as the rotation angle around the rotation axis of the initial ellipsoids from the initial into the final position.

148

Chapter 8 8.4.

Discussions

8.4.1

Effect of crystallographic variants on transition temperatures The effectiveness of martensitic block boundary over other hierarchical

martensitic boundaries in resisting dynamic brittle crack propagation has been confirmed in the current investigation, Fig. 8.1, 8.7, 8.8 and Table 8.3. During martensitic transformation, out of the 24 possible K-S variants, the adjacent variants forming block boundaries between them belong to three different Bain variant groups as categorized by Bain strain axis along γ, γ and γ [31]. In case of the variants with the same Bain compression axis, the rotation required for transformation is less [321]. On the other hand, the variants belonging to different Bain groups, that are separated by block boundaries, are highly misoriented and have large angular difference (i.e. > 30°, Fig. 8.1b) between their respective {100} cleavage planes. Thus block boundaries significantly deflect the cleavage crack propagation path (> 15°). It is worth mentioning at this point that during the observation of brittle crack path one should consider the projection angle of the actual cleavage angle between the adjacent crystals on the surface of observation. Only direct observation of the crack path [322] may often mislead in assessing the effectiveness of a particular martensitic boundary in crack deviation. The observation that in few cases the cleavage cracks propagate across the block boundaries without much angular deviation can be primarily attributed to the projection angle of the actual crack deviation on the surface of observation Fig. 8.7, 8.8 and Table 8.3. A quantitative analysis of the fractions of different variants present in the microstructures of the investigated specimens has been carried out following the procedure mentioned in Section 8.2.1 as represented by the Euler map of 875HR1100NT specimen, Fig. 8.17. Variant selection during controlled-rolling and subsequent normalization treatment (at or above TNR) can create a variation in length for certain combinations of inter-variant boundaries for different investigated samples [55,321], Fig. 8.18. Almost a uniform distribution of inter-variant boundaries can be observed for 950HR-1025NT, 1000HR-950NT, 1000HR-1025NT and 1050HR-950NT samples, Fig. 8.18(e, g, h, j). This indicates to the possibility of formation of all types of martensitic boundaries (including sub-block, block and packet boundaries), which provided ~82.6% and ~74% boundaries to have greater than 15° misorientation angle

149

Chapter 8  and cleavage angle, respectively, Fig. 8.3. As a consequence, the EGS{100} was lower and that resulted in low DBTT for these samples, Table 8.2. On the other hand, 950HR950NT, 950HR-1100NT and 1000HR-1100NT samples indicated higher length fractions of some specific inter-variant boundaries, typically V6-V20, V5-V17, V18V21, V1-V9 and V3-V15 types of packet boundaries, as well as V1-V5, V9-V10, V2V6, V8-V12, V13-V15 and V13-V18 types of block boundaries, Fig. 8.18(d, f, i). Interestingly, these boundaries apart from being high cleavage angle boundaries, also form very high twist angles (17°-29°), Fig. 8.1c, when the grain boundaries are perpendicular to the rolling direction (i.e. grain boundary plane normal//RD), which is also the direction of propagation of the primary brittle crack in the present scenario, Fig. 3.6. Higher twist angle boundaries are known to consume a significant amount of deformation energy during the crack propagation across the boundary that usually form a step on the crack path [173,305,323]. Higher length fraction of these boundaries can reduce the DBTT by providing a strong barrier to the crack propagation path irrespective of the effective grain size, Table 2 and Fig. 6. In case of 875HR-950NT, 875HR-1025NT and 1050HR-1100NT samples, although the distribution of inter-variant boundaries are not specifically restricted to be of certain types, some of these having higher length fraction (i.e. V1-V14, V7-V14, V7V14, V11-V14 and V3-V8 type of packet boundaries), Fig. 8.18(a, b, l), formed high cleavage angles and thereby reduced the EGS{100} values, Table 8.2. However, small twist angles (2°-9°) generated by these boundaries (Fig. 8.1c) possibly made these ineffective in resisting the cleavage crack propagation. Hence, the DBTT were found to be relatively higher in these specimens in spite of their small effective grain size, Table 8.2. Therefore, mere estimation of the effective grain size cannot be a true representation of the resistance offered by the hierarchical martensitic microstructure to the cleavage crack propagation and consideration of the twist angle of different boundaries in also necessary.

150

Chapter 8

Fig. 8.17: Typical Euler angle map of 875HR-1100NT sample with reconstructed prioraustenite grain structure shown in inset. The identified variants in one of the prioraustenite grain, ‘A’ (orientation: φ1=90.3°, φ=104.1° and φ2=180.8°) are indicated in the microstructure.

151

Chapter 8 

Fig. 8.18: Distributions of inter-variants boundaries for different investigated samples. (a) 875HR-950NT, (b) 875HR-1025NT, (c) 875HR-1100NT, (d) 950HR-950NT, (e) 950HR-1025NT, (f) 950HR-1100NT, (g) 1000HR-950NT, (h) 1000HR-1025NT, (i) 1000HR-1100NT (j) 1050HR-950NT, (k) 1050HR-1025NT, (l) 1050HR-1100NT.

152

Chapter 8 8.4.2

Elasticity and micro-plasticity of cleavage crack propagation The propagation of brittle ‘cleavage’ cracks in transgranular fashion across the

inter-crystalline boundaries is primarily elastic in nature. Therefore, mismatch in elastic properties due to the crystallographic anisotropy [12] as a result of intrinsic difference in orientations of different variants may impose a resistance to an advancing crack front as it propagates from one variant to another. This argument can be justified by showing the elastic modulus map of the microstructures surrounding a cleavage crack, Fig.

8.19a. In the modulus maps, the cleavage cracks exhibit significant deviation (30°-45°) at places where strong difference in local elastic modulus values are observed as evident from different colour codes in Fig. 8.19a. However, the micro-plasticity associated with the crack propagation can be another contributing factor for local variation in orientation. Dislocation emission ahead of a crack tip during cleavage fracture can cause micro-plasticity [324], which can be visualized in terms of local strain distribution as represented by the Kernel Average Misorientation (KAM) maps of the location surrounding the cleavage crack, Fig. 8.19b. As the crack front propagates, the elastic interaction between cleavage crack and emitted dislocations ahead of the crack tip can cause the crack shielding by generating a mixed mode loading condition [324]. Thus, large twist angles for certain adjacent variants (as discussed in Section 8.4.1) are expected to generate micro-plasticity and thereby, can increase the impact toughness even below the transition temperature, Fig. 8.19b.

153

Chapter 8 

Fig. 8.19: Similar secondary cleavage crack regions as that shown in Fig. 8.7. (a) Elastic modulus maps show variation in modulus (in terms of colours) at the locations of significant crack deviations, and (b) strain distributions adjacent to crack path obtained from Kernel Average Misorientation (KAM) maps. Colour legends are shown in upperleft corners. 8.4.3

Rotational phenomenon within the adiabatic shear bands Rotation of grains (discussed in Section 8.2.6) within the adiabatic shear bands

(ASBs), have been estimated using crystal plasticity simulation for the specimens impact tested at higher temperatures, Fig. 8.16. The salient features of these rotational phenomena can be attributed to the following reasons.



Smaller amount of grain rotations in small grained S1 and S2 samples can be attributed to the initial elongated grain shape along the rolling direction (i.e. direction of applied shear), Fig. 8.13(a, b), which restricted further grain rotations during shear deformation. Apart from this, higher number of active slip systems per grain (> 3.6) involving both {110} and {112} slip planes in small grained specimens (S1 and S2), could have uniformly accommodated the plastic strain without substantial amount of grain rotation, Fig. 8.20, as multiple slip causes less change in grain orientation as opposed to single slip (Boas and Hargreaves, 1948). Moreover, geometrically necessary dislocations favour the uniform deformation in very fine-grained polycrystal sample, and result in Taylor mode of interaction between the grains under deformation [318–320]. This could be the reason why ‘Secant model’, which is nearer to the Taylor 154

Chapter 8 deformation mode, provided satisfactory prediction of deformation texture in S1 sample having smallest grain size and grain size spread among the investigated specimens.



Severe amount of grain rotation in large grained specimens (L1 and L2) can be attributed to the following factors: (i) non-uniform constraint effect imposed by the adjacent grains on the deforming grain and (ii) non-uniform deformation between the different regions within a grain (local strain is expected to be higher at the vicinity of grain boundary than at grain interior, [49]). As a result, the locations under higher constraint will require substantial amount of slip activity from the conjugate slip systems (i.e. {112} type), apart from the involvement of the primary ones (i.e. {110} type). This phenomenon has been observed at different stages along the strain path of the L1 and L2 samples, where the relative contributions of the slip systems to the deformation process were changed and sometimes slip activity of {112} type of slip systems surpass that of {110} type of slip systems, Fig. 8.21(c and d). However, the relative contribution of different slip systems to the shear deformation remains almost constant in S1 and S2 samples, Fig. 8.21(a-b). Predominance of {112} slip systems can enhance the grain rotation in order to align the slip directions with that of the loading axis to keep the resolved shear stress sufficient for further deformation [326], as schematically shown in Fig. 8.22. Moreover, lower number of average active slip systems per grain in L1 specimen might have resulted in higher rotation of grain in order to continue the shear deformation. However, non-uniform deformation effect might be there at a small extent even in S2 specimen, due to the presence of larger grains than S1, as a result of which ‘Tangent model’ offered satisfactory texture simulation in S2 along with L1.



Very large fraction of area (~ 85%) that is covered by grains larger than the average grain size provides regions of non-uniform strain accommodation along the strain path. Different regions within the large grains are subjected to different amount of plastic deformation depending upon the directionality of the slip systems with respect to the loading axis. This phenomenon also indicates that significant plastic strain gradient arises within the large prioraustenite grains as the area fraction and size of those grains exceed a threshold

155

Chapter 8  level. This could be the reason why relative directional compliance scheme [316] provided a satisfactory texture prediction in L2.

Fig. 8.20: Number of active slip systems per grain within ASB regions of different investigated samples along the strain path.

Fig. 8.21: Fraction of slip activity for {110} and {112} slip systems along the strain path within the ASB regions of (a) S1, (b) S2, (c) L1 and (d) L2 samples.

156

Chapter 8

Fig. 8.22: Schematic representation of a single crystal subjected to loading. The crystal initially deforms primarily with slip primarily in slip system, S1. Lattice rotations (α) occurs such that slip system S2 rotates toward the loading axis which eventually leads to the conjugate system to be active in S2` state. Therefore, slip activity in conjugate slip systems promotes lattice rotation. 8.4.4

Microstructural damage within adiabatic shear bands Different grain domains within a shear band exhibited intra-grain misorientation

or deviation from the average orientation, as shown by the grain orientation spread (GOS) maps of all the investigated samples, Fig. 8.23. Deformation strain gradient and the rigid body rotation associated with the simple shear loading configuration result in a variation in crystallographic orientation (or orientation perturbation) within a grain [327]. Larger PAGS and significant amount of lattice rotation contribute to higher orientation spread in L1 and L2 specimens as compared to S1 and S2 specimens, Fig.

8.23. Additionally, the variation in deformation strain gradient inside the very large prior-austenite grains in L2 results in highest orientation spread (~4.7) within its shear bands amongst the other investigated samples, Fig. 8.23. Large grained L1 and L2 specimens show extensive amount of grain fragmentation within the ASBs, as shown by polycrystalline diffraction patterns and corresponding dark field images taken from the samples prepared by FIB technique,

Fig. 8.24(a and b). Similar fragmentation of grains into nanocrystalline size range was also observed in austenitic stainless steels when subjected to high strain-rate deformation [284,298]. Significant deviation in orientation over different domains within a grain of L1 (~ 3.8) and L2 (~ 4.7) samples (Fig. 8.24) are expected to develop deformation induced incidental dislocation boundaries (IDBs) and geometrically

157

Chapter 8  necessary boundaries (GNBs). These GNBs separate the regions undergoing independent deformation owing to significant lattice rotation. This results in grain subdivision by creating high misorientation angle around these regions [328]. However, smaller grain size and negligible rotational effect during deformation occasionally result in grain fragmentation in the S1 and S2 specimens, Fig. 8.24(c and d). In case of these samples, mostly deformed laths containing extensive amount of dislocations and well defined sub-cells are primarily seen, Fig. 8.24(c and d). However, in some cases, dislocations crossing the grain boundaries were observed in S1 specimen, Fig. 8.24e. This phenomenon can be attributed to the availability of a large number of active slip systems (~ 6) within each grain of S1 specimen, similar to the observation of Ogilvie, where intersecting slip planes at the grain boundary of a fcc metal created a small angular deviation (2º) for different slip directions [329].

Fig. 8.23: Intra-grain orientation deviations within adiabatic shear band regions are shown by Grain orientation spread (GOS) maps of (a) S1, (b) S2, (c) L1 and (d) L2 samples.

158

Chapter 8

Fig. 8.24: Typical transmission micrographs from ASB regions of (a) L2 and (b) L1 showing fragmented grains in bright field (BF) and dark field (DF) mode with selected area diffraction (SAD) pattern in inset, and of (c) L2 and (d) S1 indicating rarely visible fragmented grains, sub-cells and high dislocation density within laths shown by yellow, blue and red arrows, respectively. (e) Examples of dislocations crossing GBs in S1 sample, higher magnification BF image of A and B zones from (d). Interestingly, high resolution electron microscopic (HREM) study revealed the presence of partial disclinations adjacent to the localized cracks, in the order of few nanometer, situated mostly at the grain boundary junctions of L2 sample, Fig. 8.25a. These disclinations (analogous to dislocation but cause defect in rotational symmetry [330]) facilitate the grain rotation and terminate at tilt grain boundaries during plastic deformation, Fig. 8.25a. Disclination dipoles of {110}, shown in Fig. 8.25a3, creates an angular rotation of 12º with their parallel counterpart, which is relatively higher as

159

Chapter 8  compared to the earlier report on nanocrystalline Fe [331], without the formation of any crack. The partial disclination however, forms micro-crack (~ 2 nm) by the coalescence of small voids, as indicated by yellow arrow in Fig. 8.25a2-a3. Although disclinations were found in some places within the ASBs of L1 sample, these were found as disclination dipoles without the formation of cracks or voids, Fig. 8.25b. Variation in rotation due to slip, and severe elastic-plastic incompatibility exists as a result of strain gradient within the large prior-austenite grains in L2 specimen might be the reason behind the formation of partial disclinations. Elastic stress of partial disclination leads to the formation of micro-cracks, whilst these stresses get nullified in case of disclination dipoles. However, disclination phenomenon was not found in S1 and S2 specimens possibly due to negligible amount of rotational effects and relatively narrow grain size spread.

Fig. 8.25: (a) Typical HREM image of micro-voids in the formation of crack near GB triple points in ASB region of L2. Enlarged view shows disclination dipoles and partial disclination indicated by red and yellow ovals, respectively, (b) presence of similar disclination dipoles in L1, but without formation of crack or void. Nevertheless, among the four selected samples, the DBTT of S1 and L2 were much lower than that of S2 and L1, Fig. 8.4 and Table 8.2. Therefore, neither PAGS nor grain size distribution spread are the controlling factors for DBTT. However, DBTT

160

Chapter 8 clearly decreases (which is beneficial) with the refinement in the average size of martensitic blocks from L1 (5.2 µm), S2 (4.3 µm), L2 (1.5 µm) and finally to S1 (1.1 µm) samples, Table 8.2. Thus it becomes evident that martensitic blocks act as the ‘effective grain’ in resisting the cleavage crack propagation as discussed in the previous sections. However, the lowest USE in conjunction with the higher DBTT of S2 (875HR-1100NT) resulted in the slant nature of impact transition curve, Fig. 8.4. In contrast, combination of higher USE and lower DBTT contributed to steepest transition behaviour of S1 (1000HR-950NT), followed by L2 (1000HR-1100NT) and L1 (1050HR-1100NT), Fig. 8.4. Hence, small prior-austenitic grain structures having small martensitic blocks can provide an excellent combination for increasing USE and reducing DBTT [332]. However, it is worth mentioning that apart from the grain structure precipitates can play a significant role on the impact fracture behaviour of a particle containing system irrespective of the test temperature. In the investigated steel, at low test temperatures, large Cr23C6 particles (Fig. 8.26a-b) having low fracture stress may act as the potential sites for cleavage crack initiation either by the decohesion of particlematrix interface or by the fracture of particles [12]. On other side, at high test temperatures, Cr23C6 particles are expected to nucleate large primary voids, whereas, MX type of microalloy precipitates (Fig. 8.26b-c) can be responsible in joining the primary voids by forming small secondary voids at the ligament [211]. Therefore, decreasing the volume fraction of large Cr23C6 carbides and increasing the inter-particle spacing of MX precipitates, can be beneficial from impact toughness point of view. However, that may impose an adverse effect of the elevated temperature creep resistance, which is the prime requirement of 9Cr-1Mo steel. These aspects will be studied in future.

Fig. 8.26: Typical transmission electron micrographs showing of presence of M23C6 in (a) HAADF-STEM image and (b) dark-field image, and MX type precipitates can be 161

Chapter 8  seen in (b) dark-field image as well as (c) high-magnification bright field image. M23C6 and MX precipitates are indicated by red and yellow arrows, respectively.

8.5.

Summary A modified 9Cr-1Mo steel was hot-rolled at different temperatures (1050-

875C) followed by normalized using different austenitizing temperatures (1100950C) and finally tempered at 750C. These samples having tempered martensitic microstructures were impact tested over the temperature range of +80C to -196C. The investigation of dynamic fracture at low temperature leads to the following conclusions:

 the hierarchical martensitic boundaries should be characterized in terms of angular deviation between the cleavage planes (i.e. {100} type crystallographic planes in bcc ferritic / martensitic steels) of the adjacent martensitic variants, as opposed to the misorientation angle (angle-axis) in order to assess their ability to resist the cleavage crack propagation.

 Martensitic blocks are found to be the ‘effective grain’ for low-temperature impact toughness. Block boundaries are more effective in cleavage crack retardation as compared to the packet boundaries, as considering cleavage angle across the martensitic variants all the block boundaries and ~75% of packet boundaries are the high-angle boundaries.

 Twist angle component of the martensitic boundaries that are perpendicular to the crack path is necessary to be considered as presence of higher fraction of intervariant boundaries with high-twist angle in a microstructure provides higher lowtemperature impact toughness. Among the all processed samples, four specimens: two with small average prioraustenite grain size (13-27 µm), S1 and S2, and two with large average prioraustenite grain size (~ 60 µm), L1 and L2 are selected to study the high-temperature ductile fracture mechanism under dynamic condition. The grain size spread in S2 is higher than S1, whilst the grain size spread in L2 is higher than L1. The results obtained in this investigation are summarized below:

 Shear cracking occurs along the adiabatic shear bands (ASBs) and the deformation mechanisms within the ASBs depend on the prior-austenite grain size (PAGS) and grain size spread in the microstructure.

162

Chapter 8  Samples with large PAGS (L1 and L2) show rotated ASBs indicating substantial grain rotation, which leads to significant deviation in intra-grain orientation within an ASB, and results in the grain fragmentation during plastic deformation. These can be attributed to non-uniform constraint effect imposed by the adjacent grains that activate greater slip activity at certain locations by the involvement of conjugate slip systems (i.e. {112} type), whilst, primary slip systems (i.e. {110} type) dictate the deformation at the other locations within a grain.

 Lesser extent of grain rotation in smaller grained specimens (S1 and S2) can be attributed to initial elongated PAG structure along the direction of shear (i.e. rolling direction), and higher number of active slip systems per grain (> 3.6).

 Higher amount of grain rotation in large grained specimens (L1 and L2) can be attributed to the mixed grain structure with very large spread in the grain size distribution and the associated non-uniform deformation effect within the large grains and more uniform deformation of the smaller grains.

 The sub-microscopic damage in L2 specimen occurs through the formation of partial disclinations within the large grains. Elastic stress of partial disclination can lead to the formation of micro-cracks, whilst these stresses get nullified in case of disclination dipoles (in L1 sample). This phenomenon is rarely visible in specimens having small and uniform prior-austenite grain size possibly due to the negligible amount of rotational effects.

Appendix:

Fig. 8.A1: Shear stress-shear strain curves used for determination of Voce hardening parameters. Line represent experimental data, while points are simulated ones. 163

 

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164

 

Chapter 9 Effect of normalization treatment on creep strength at different creep temperatures

 

 

Chapter 9 9.1

Introduction and objective Modified 9Cr-1Mo steel is widely used in the power generation industry, where

the material requires prolonged service at temperatures upto ~600°C [333]. Addition of V and Nb in modified 9Cr-1Mo steel is aimed to improve its creep properties as these microalloying elements form finely dispersed and stable MX types of carbides or carbonitride precipitates. These precipitates form along the martensitic lath boundaries during tempering treatment of air-hardened martensitic structure [149,150]. MX precipitates prevent lath coarsening and the resultant structural softening upon prolonged thermal exposure, which improves the high temperature creep properties of modified 9Cr-1Mo steel [157,158]. Previous studies on creep behaviour of Cr-Mo steels have extensively dealt with the mechanisms of creep deformation depending upon different operating temperatures and stress regimes [115–126]. Several studies were also carried out to understand the effects of grain structure [127,128], evolution and stability of different types of carbide (i.e. M23C6 or MX types) particles [113,129–131], the Laves phases during prolonged thermal exposure [126,132–134], and the role of special boundaries on the creep behaviour [135]. Numerous studies on 9Cr-1Mo steel investigated the evolution of microstructure in terms of formation of subgrains, and progressive changes in the nature and intensity of these boundaries, variation in dislocation density as well as formation and coarsening of different types of precipitates when subjected to thermal exposure under stress [113,136–141]. In order to use in ultra-supercritical (USC) power plants or in fuel tubes of fast breeder nuclear reactors, heat resistant modified 9Cr-1Mo steel is thermomechanically processed (hot-forged and/or hot-rolled) and finally used in ‘normalized

and

tempered’

heat-treated

condition

[18,20,28,118,142–147].

Normalization and subsequent tempering treatment change the martensitic microstructures, the precipitate particles (nature, size, fraction and distribution), and even the crystallographic texture of 9Cr-1Mo steel as discussed in Chapter 6 and also reported earlier [19,148]. It is already shown in Chapter 6 that the normalization treatment can have a significant influence on the microstructure, tensile and Charpy impact properties of the investigated steel. Therefore, the effect of normalization temperature on the creep properties needs to be studied in order to optimize the processing parameters for achieving the desired creep resistance in modified 9Cr-1Mo steel under service condition. In this chapter, the results on the study of the effect of 165

Chapter 9  normalization temperature on the creep properties and the evolution of microstructures and crystallographic texture after creep exposure (at different temperatures under constant stress level) have been discussed.

9.2. Results and Discussion 9.2.1

Creep behaviour of investigated samples Creep testing was carried out at three different temperatures, 550 ºC, 600 ºC and

650 ºC, Fig. 9.1. The creep tested specimens are denoted by their normalization temperature (NT), followed by creep testing temperature (e.g. 950NT-550 suggests normalization at 950 ºC followed by tempering, and creep testing at 550 ºC). The creeprupture behaviour of all the investigated samples are represented in creep strain vs. time to rupture plot as shown in Fig. 9.2a. These plots primarily represent typical creep curves with three different regions of creep deformation: primary stage, secondary or steady state and tertiary creep regime. During low temperature creep test (i.e. 550 ºC), only 950NT-550 specimen was found to be ruptured within the duration allowed for testing, whilst 1025NT-550 and 1100NT-550 samples reached a steady state after creep test for ~1813 h and ~1110 h, respectively, Fig. 9.2b. On the other hand, the maximum time to rupture at 600 ºC has been recorded for 1100NT-600 sample followed by 1025NT-600 and 950NT-600 specimens, whilst the strain to rupture is found to be the highest in 1025NT-600 sample, Fig. 9.2c. At high test temperature (i.e. 650 ºC), interestingly, the 1025NT-650 sample has shown the maximum time to rupture as well as the highest creep strain to rupture as compared to the other samples, Fig. 9.2d. The minimum rupture life and creep strain are found in case of 950NT-650 and 1100NT650 samples, respectively, Fig. 9.2d.

Fig. 9.1: Schematic diagram showing normalized sample subjected to three different creep test temperatures.

166

Chapter 9 All the investigated samples indicate increase in minimum creep rate with increase in respective creep test temperature, Fig. 9.2e. At low testing temperature (i.e. 550 ºC) regime, the 1025NT-550 has shown the slowest minimum creep strain rate (order of 10-12 s-1) followed by 950NT-550 and 1100NT-550, Fig. 9.2e. Although the minimum creep rate exhibited by the 950NT-550 sample is lower than that by the 1100NT-550, the former test sample has not shown better creep life than the latter, Fig.

9.2(b,e). While 950NT-550 was ruptured after 620 hours, the 1100NT-550 reached steady state creep regime after 1200 hours, Fig. 9.2b. This phenomenon can be attributed to the absence of clearly distinguishable steady state regime for 950NT-550 sample during creep deformation, Fig. 9.2b. Therefore, analysing the creep-resistant ability of any material based on minimum creep rate may sometime mislead, rather time to rupture should also be considered in this regard, Fig. 9.2f. However, at intermediate test temperature (i.e. 600 ºC), the minimum creep rates of 950NT-600 and 1025NT-600 are significantly higher as compared to those creep tested at 550 ºC, whilst the 1100NT600 has the lowest value, Fig. 9.2e. Similar to the 950NT-550, 1100NT-650 has shown the lowest creep rate, although it has relatively lower time to rupture as compared to the 1025NT-650, Fig. 9.2e. Throughout the test temperature regime (i.e. 550ºC-650ºC) 1100NT showed marginal increase, whilst 1025NT and 950NT indicated significantly jump in their minimum creep rate with the increase in creep test temperature. All the creep properties of the normalized samples are listed in Table 9.1.

Fig. 9.2: (a) Creep strain vs. time to rupture plots of all the investigated samples. Creep curves of normalized samples tested at 550 °C (b), 600 °C (c) and 650 °C (d). Minimum creep rate and duration of creep tests of the investigated samples are shown in (e) and (f), respectively. 167

Chapter 9  Table 9.1: The creep properties of the normalized samples. Creep test temperature

Normalizing Min Creep Time to temperature rate (s-1) rupture (hrs)

950°C 1.60E-09 620.4 1025°C 8.30E-12 2037.7* 550°C 1100°C 3.30E-08 1285.7* 950°C 3.60E-05 9.1 1025°C 1.85E-05 41.4 600°C 1100°C 1.67E-07 323.2 950°C 4.00E-03 0.09 1025°C 1.24E-04 36.7 650°C 1100°C 6.67E-07 9.3 * Not ruptured up to the duration of creep test

Strain to failure (%) 3.048 0.81 0.42 2.463 3.565 1.45 5.69 16.34 0.8

9.2.2. Effect and evolution of microstructure in relation to creep On being subjected to at low temperature (550°C) creep test, 950NT-550 had the lowest creep life, whilst the 1025NT-550 reached the steady state creep regime at a much later stage with the minimum creep rate being the lowest among all, Fig. 9.2(b,e). Significantly higher amount of strain free regions were found in the microstructure of ruptured 950NT-550 sample, as shown in the grain orientation spread (GOS) maps (using GOS < 1) in Fig. 9.3a. However, the sample after normalization at 950 °C had high dislocation density with negligible strain free regions, Fig. 9.4a. The reason may be that the normalization temperature of 950 °C is very close to the recrystallization stop temperature of the investigated steel and sufficient structural softening (by dislocation rearrangement) could not take place at that temperature. Therefore, the strain free regions as observed in the creep tested 950NT-550 sample are formed by lath coarsening of highly strained grains. Similar formation of equiaxed strain-free grains during creep of 9Cr-1Mo have also been reported in earlier works [131,137]. These strain free grains have an adverse effect on the creep resistance [131] of 950NT550 and lowers the time to rupture. In a similar manner, the presence of a large amount of strain free grains is expected to have increased the minimum creep rate of 1100NT550, Fig. 9.3c, 9.4a and 9.2e. However, significant amount of these strain free grains were already there after the normalization treatment at high temperature (1100ºC), apart from formation of some amount of these grains during the creep.

168

Chapter 9

Fig. 9.3:Grain orientation spread (GOS) maps of (a) 950NT-550, (b) 1025NT-550, (c) 1100NT-550, (d) 950NT-600, (e) 1025NT-600, (f) 1100NT-600, (g) 950NT-650, (h) 1025NT-650 and (i) 1100NT-650 samples, where GOS < 1 (blue in color) indicates the recrystallized grains. .

Fig. 9.4: (a) Area fraction of strain-free regions in the microstructures and (b) ratio of primary creep strain to total creep strains of the investigated normalized samples as function of creep test temperatures.

169

Chapter 9  Less amount of strain free regions were present in 1100NT-600 creep tested sample, which is consistent with its lower creep strain rate and longer time to rupture, Fig. 9.3f, 9.4a and 9.2e. Higher ratio of primary creep strain to total rupture strain in 1025NT-600 as compared to 1100NT-600 (Fig. 9.4b) is expected to have promoted dynamic recovery during the later stage of creep in 1025NT-600, (please see Section 9.2.3), which increased the fraction of strain free grains and reduced the time to rupture as compared to 1100NT-600, Fig. 9.3(e,f), 9.4a, Fig. 9.2f. However, at very high temperature (650ºC), both 950NT-650 and 1100NT-650 samples showed rapid increase in the fraction of strain free regions (Fig. 9.4a) as a result of dynamic recovery being promoted by significant amount of strain hardening as shown by the large ratio of primary creep strain to rupture strain, Fig. 9.4b. This caused a significant reduction in creep life of these specimens, Fig. 9.2f. In contrast, the 1025NT-650 has shown a steady moderate growth in fraction of strain free regions in the microstructure due to the decrease in contribution of strain hardening before steady state creep regime, which contributed to only a marginal decrease in rupture life of the 1025NT-650 sample as compared to that of the 1025NT-600, Fig. 9.4, Fig. 9.3(h,i) and Fig. 9.2f . In the sample normalized at 950°C, the microstructural features can be summarized as follows: (i) prior-austenite grain sizes and the martensite packet sizes were fine, (ii) intensities of dislocation substructure and low-angle ‘sub-boundaries’ were very high, and (iii) normalization temperature of 950°C was not sufficient to dissolve the pre-existing precipitates and hence, coarse precipitates remained in the microstructure. During creep exposure, dislocation substructure could have increased the diffusivity (by allowing dislocation pipe diffusion), which coarsened the precipitates further, and as a result, the precipitate pining effect at the martensitic lath boundaries decreased. Presence of sufficient prior-austenite grain boundaries (due to the finer microstructure) and the martensitic packet boundaries further accelerated the diffusion process significantly. This possibly resulted in lath coarsening and helped in the formation of strain free regions in the microstructures. Basically, an increase in the normalization temperature (for 1025NT and 1100NT specimens) resulted in (i) increase in both prior-austenite grain size and martensitic packet size, (ii) sufficiently coarsened dislocation substructure, and (iii) opportunity for the dissociation of pre-existing precipitates during cooling / tempering. All these factors are expected to slow down the diffusion controlled processes, and also

170

Chapter 9 stabilize the microstructure during creep exposure, resulting in an improvement in creep life and decrease in minimum creep rate. SEM micrographs of the transverse cross-section just below the surfaces of the creep tested specimens are presented in Fig. 9.5. In the specimens creep-ruptured at 550°C, predominantly slip line markings can be seen, which possibly indicate dislocation glide is the dominant creep deformation mechanism at 550ºC, Fig. 9.5(b-

c). Interestingly, certain location of the specimens tested at 600ºC could have also undergone the grain-boundary sliding phenomenon as indicated by presence of intergranular cracks depicted in Fig. 9.5(d-f). In high temperature (650ºC) creepdeformed specimens, formation of creep cavities around the coarse particle-matrix interfaces are the predominant mechanism of creep damage formation within the microstructures, Fig. 9.5(g-i).

Fig. 9.5: Sub-surface creep damage evolution near ruptured surfaces of maps of (a) 950NT-550, (b) 1025NT-550, (c) 1100NT-550, (d) 950NT-600, (e) 1025NT-600, (f) 1100NT-600, (g) 950NT-650, (h) 1025NT-650 and (i) 1100NT-650 samples. Slip markings in (b) and (d), grain boundary slidings in (d-f), and creep cavities in (g-i) are indicated by arrows. 9.2.3

Effect and evolution of precipitates in relation to creep Large Cr23C6 carbides and smaller sized Nb(C,N) precipitates remain mostly

undissolved during low temperature normalization treatment at 950C, Fig. 6.8. These

171

Chapter 9  precipitates get coarsened during normalization and tempering treatments. On the other hand, the highest normalization temperature of 1100C is sufficient to dissolve almost all of the pre-existing precipitates as predicted thermodynamically (Fig. 6.8), and subsequent tempering treatment can freshly form fine Cr23C6 and MX carbonitrides, Table 6.1. During normalization at intermediate temperature of 1025C, a portion of pre-existing Cr23C6 and Nb(C,N) precipitates are expected to be present, whereas the remaining portion is dissolved in austenite (Fig. 6.8). As a result, a combination of preexisting and newly formed precipitates was present in normalized and tempered condition. Therefore, larger inter-particle spacing in 950NT samples as compared to 1025NT and 1100NT specimens contributes to the higher creep rate in the former in comparison to the later ones following Orowan’s equation:

  b





 b     tc  tg tc

Eqn. 9.1

where,  denotes the creep rate. ρ, b and λ are dislocation density, Burgers vector and inter-particle spacing, respectively. tc and tg are the time taken by the dislocations to climb over a particle and to glide, respectively. Being tg