Doctoral Thesis in Materials Science and Engineering
Thermomechanical Processing of
Nickel-Base Alloy 825
MUNIR AL-SAADI
Thermomechanical Processing of
Nickel-Base Alloy 825
MUNIR AL-SAADI
Doctoral Thesis in Materials Science and Engineering KTH Royal Institute of Technology
Stockholm, Sweden 2021
Academic Dissertation which, with due permission of the KTH Royal Institute of Technology, is submitted for public defence for the Degree of Doctor of Technology on Monday the 14th June 2021, at 10 a.m. Digital and Green room, Osquars backe 31, Södra tornet plan 4, Stockholm.
© Munir Al-Saadi
© Munir Al-Saadi, Fredrik Sandberg, Andrey Karasev, Stefan Jonsson, Pär Jönsson, Supplement I
© Munir Al-Saadi, Wangzhong Mu, Christopher N. Hulme-Smith, Fredrik Sandberg, Pär Jönsson, Supplement II © Munir Al-Saadi, Christopher N. Hulme-Smith, Fredrik Sandberg, Pär Jönsson, Supplement III
© Munir Al-Saadi, Christopher N. Hulme-Smith, Fredrik Sandberg, Pär Jönsson, Supplement IV © Munir Al-Saadi, Christopher N. Hulme-Smith, Fredrik Sandberg, Pär Jönsson, Supplement V TRITA-ITM-AVL 2021:32
ISBN: 978-91-7873-827-4
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4 Abstract
In this thesis, the dynamic recrystallization behaviour of a cast Alloy 825 material was studied using a Gleeble-3800 thermo-simulator by performing single-hot compression experiments. Optical microscopy and electron backscatter diffraction were utilized to characterize the microstructure. Dynamic recrystallization is not considerable in the as-cast alloy and dislocation recovery is deemed to be dominant. Based on this finding, the effect of adding trace amounts of alloying additions on the mechanical properties of cast alloy 825 was studied, with emphasis on whether or not dynamic recrystallization occurred. The results show that dynamic recrystallization was more prevalent under all test conditions in samples containing a trace amount of magnesium, but not for the conventional alloy. However, alloying with trace magnesium did not lead to an improvement of the mechanical properties. Instead, processing maps for hot forging of conventional Alloy 825 were required to identify optimal working parameters and to achieve dynamic recrystallization. The hot deformation behavior of cast Alloy 825 was characterized by using dynamic materials modelling of hot compression data. The results show that the maximum power dissipation efficiency is over 35%. The highest efficiency is
achieved in the temperature range of 1100 ℃ - 1250 ℃ and in
strain rates in the range of 0.01 ≤ / s ≤ 0.1. The optimum
processing parameters for good strain hardening are obtained in
the temperature range between 950 ℃ and 1100 ℃ with strain
rates of 0.3 ≤ / s ≤ 10.0. In addition, the influence of the
deformation level on the recrystallization and microstructural changes in Alloy 825 during hot forging operations at temperatures between 950 °C and 1200 °C was studied. The maximum yield strength and ultimate tensile strength were obtained after forging to achieve a true strain of 0.9 were 413 MPa and 622 MPa, respectively, with a ductility of 40%.
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However, Alloy 825 is often supplied as annealed bars. Therefore, the effect of the forging strain magnitude and subsequent annealing on the microstructure, strengthening mechanisms and room temperature mechanical properties were investigated to assess the suitability of current industrial practice. The results showed that the majority of strengthening was attributed to grain refinement, the dislocation densities that arise due to the large forging strain, and due to solid solution strengthening. The results of calculations are in excellent agreement with experimental data, with less than 1% difference. These results can be used by future researchers and industry to predict the strength of Alloy 825 and similar alloys, especially in material after a completed hot forging operation.
Keywords: Alloy 825; Cast and wrought structures; Hot compression test; Microstructural evaluation, Modelling, Yield strength, Strengthening mechanisms.
6 Sammanfattning
I denna avhandling studerades effekten av deformationsgrad på dynamisk rekristallisation av gjuten Alloy 825 genom experiment i Gleeble-3800 termomekanisk simulator. Målet med arbetet är att använda ljusmikroskop och EBSD för att studera dynamisk rekristallisation i A 825 och dess effekt på mikrostrukturen. Dynamisk rekristallisation är inte betydande, men istället har en substruktur formats med lågvinkliga korngränser.
Baserat på dessa resultat, studerades effekten av små magnesiumtillsatser på de mekaniska egenskaperna av gjuten struktur med fokus på huruvida dynamisk rekristallisation sker eller inte. Resultaten visar att dynamisk rekristallisation var vanligare under alla testförhållanden i prover innehållande magnesium, men inte i något fall av konventionell A 825. Legering med magnesium ledde emellertid inte till en förbättring av de mekaniska egenskaperna. Av denna anledning drog man slutsatsen att en varmbearbetningskarta för smitt konventionell (magnesiumfri) A 825 krävdes för att identifiera
optimala processparametrar och uppnå dynamisk
rekristallisation.
Deformationen av gjuten A 825 undersöktes med hjälp av en dynamisk materialmodell och data från kompressionsprover. Resultaten visar att den maximala effektöverföringen är över 35 %. Den högsta verkningsgraden är vid ett temperaturintervall av
1100 ℃ - 1250 ℃ och en töjningshastighet på 0.01 ≤ / s ≤
0.1 . De optimala varmbearbetningsparametrarna för god deformationshärdning erhålls i temperaturområdet mellan 950
℃ och 1100 °C med en töjningshastighet av 0.3 ≤ / s ≤ 10.0.
Vidare undersöktes effekten av reduktionsgrad på
rekristallisation och mikrostrukturutveckling vid smide inom temperaturområde 950° C och 1200° C. Den maximala sträck- och brottgränsen erhölls efter smide till sann töjning av 0,9.
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Sträckgränsen var 413 MPa och brottgränsen 622 MPa, med en brottförlängning på 40 %. Emellertid levereras materialet ofta i släckglödgat tillstånd. Därför undersöktes effekten av smide med olika reduktionsgrad följt av mjukglödgning där härdningsmekanismer och mekaniska egenskaper vid rumstemperatur undersöktes. Detta genomfördes för att bedöma lämpligheten av nuvarande industriell praxis. Resultaten visade att majoriteten av hårdnade tillskrevs små korn, dislokationstäthet som uppstod på grund av den stora smides-deformationen och härdning genom fast lösning.
Resultaten av beräkningarna överensstämmer med
experimentella data med mindre än 1 % skillnad. Resultaten kan i framtiden användas av andra forskare och i industrin för optimering av mekaniska egenskaperna för A 825 och liknade legeringar.
Nyckelorder: Legering 825; Gjutet och bearbetat strukturer; Varm kompression test; Microstrukturutveckling, Modellering, Sträckgräns, Härdning mekanismer
8 Acknowledgements
First of all, I should thank Professor Olle Wijk (previous R&D Chef), who recruited me 2007 at R&D Sandvik AB when I drove him by Taxi. I must in this acknowledgement take the opportunity to thank colleagues at R&D Sandvik AB. Anders Bergman for his help in performing the Gleeble tests and in various technical matters related with the Gleeble test equipment. I must thank Jerry Lindqvist and Lisa Lautrup for all their help in performing the FIB and TEM tests. I am grateful to Dr. Raveendra Siriki for his constructive suggestion and technical support throughout EBSD course of this investigation. Pär Hedqvist and Mats Kihlander are appreciated for help with metallographic preparations. The help from Emil Cederberg and Mats Kihlander with performing microhardness test is greatly acknowledged. I would like to thank Fredrik Sandberg, Olle Sundqvist, Jan Haraldsson, Karin Antonsson, Mattias Gärdsback, Petter A Persson and Christina Haraldsson for their valuable suggestion and guidance on carrying out this research. I wish to express my deep gratitude to Sandvik Material Technology for their grant which made this study possible and permission to publish this work.
I am very grateful to my supervisor Assistant Professor Christopher Hulme for all his enlightening comments and discussions that we had and all he taught me with his knowledge and patience. Many thanks to my other supervisors Pär Jönsson, Andrey Karasev, and Wangzhong Mu for their support, guidance and encouragement during this project.
Finally, I have to express my deepest gratitude to dear Jinan, Siradj, Hanna for all their continual encouragement, patient and support.
9 List of appended supplements
The following supplements constitute the basis of this thesis.
Supplement I: Microstructure characterisation in alloy 825.
Munir Al-Saadi, Fredrik Sandberg, Andrey Karasev; Stefan Jonsson; and Pär G. Jönsson.
Procedia Manufacturing, 2018, 15, 1626–1634.
Supplement II: Effect of Trace Magnesium Additions on the Dynamic Recrystallization in Cast Alloy 825 after One-Hit Hot-Deformation.
Munir Al-Saadi, Wangzhong Mu, Christopher N. Hulme-Smith, Fredrik Sandberg and Pär G. Jönsson
Metals, 11(1), 36 (2020)
Supplement III: Hot Deformation Behaviour and Processing Map of Cast Alloy 825
Munir Al-Saadi, Fredrik Sandberg, Christopher N. Hulme-Smith, Pär Jönsson. Submitted for publication in Journal of Materials Engineering and Performance
Supplement IV: Influence of Strain Magnitude on Microstructure, Texture and Mechanical Properties of Alloy 825 during hot-forging
Munir Al-Saadi, Fredrik Sandberg, Christopher N. Hulme-Smith, Pär Jönsson. Submitted for publication in Metallurgical and Material Transaction A
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Supplement V: Modelling of Strengthening
Mechanisms in Wrought Nickel-Based 825 Alloy Subjected to Solution Annealing
Munir Al-Saadi, Fredrik Sandberg, Christopher N. Hulme-Smith, Pär Jönsson. Submitted for publication in Journal of Metals and Materials International
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The contributions by the author of this thesis to the above supplements are as follows:
Supplement I: Performed all of the literature survey, evaluation and major part of writing. Performed all EBSD scans, observations, and analyses. Handled, studied, and evaluated all data from the experimental work: single-hit compression tests and hardness measurements. Performed all of the observations, and analyses of optical microscopy. Wrote major part of the oral presentation and participated in the conference.
Supplement II: Performed all of the literature survey, evaluation and major part of writing. Performed all EBSD scans, observations, and analyses. Handled, studied, evaluated, and modelled all the data from the experimental single-hit compression tests and hardness measurements. Performed all of the sample preparation, micro-etching, observations, and analyses of optical microscopy. Performed EBSD and FEG-SEM sample preparation, observations and analyses of SEM-EDS work.
Supplement III: Performed all of the literature survey, evaluation and major part of writing. Performed all EBSD scans, observations, and analyses. Handled, studied, evaluated, and analysed all the data from experimental work single-hit compression tests and hardness measurements. Performed all of the observations, and analyses of optical microscopy. Performed all simulation of the hardness and power dissipation maps. Performed all FEG-SEM scans, observations and analyses of SEM-EDS work.
Supplement IV: Performed all of the literature survey, evaluation and major part of writing. Performed all EBSD scans, TSL OIM Analyzer software, observations, and analyses. Performed all of the experimental work at forging mill (shop). Performed of EBSD and FEG-SEM sample preparations, observations and analyses of EBSD and SEM-EDS work.
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Handled, studied, and evaluation all the data from experimental work tensile testing, hardness testing, and TEM. Performed of EBSD and FEG-SEM sample preparations, observations and analyses using EBSD and SEM-EDS. Did not performed the experimental tensile testing and TEM scan. Did not performed the Gleeble tests, micro-hardness tests and major part of EBSD sample preparation.
Supplement V: Performed all of the literature survey, evaluation and major part of the writing. Performed all EBSD scans, observations, and analyses. Performed EBSD and FEG-SEM sample preparation, observations and analyses of EBSD and SEM-EDS work. Performed all thermodynamic calculations of the equilibrium phase fractions. Handled, studied, and evaluated all the data from experimental tensile testing and Powder X-ray Diffraction.
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Other Supplements not included in the thesis:
• Comparative Study of Microstructures Evolution
of Columnar and Equiaxed Grain Structures in Alloy 825 after Hot Compression, Al-Saadi, M.; Sandberg, F.; Karasev, A.; Jönsson, P. In Proceedings of the 3rd International Conference on Ingot Casting, Rolling and Forging, ICRF2018, Stockholm, Sweden, 16– 19 October 2018; pp. 107–115.
• A study of the static recrystallization behaviour of
cast Alloy 825 after hot-compressions, Al-Saadi, M.; Sandberg, F.; Hulme-Smith, C.; Karasev, A.; Jönsson, P.G. J. Phys. Conf. Ser. 2019, 1270, 012023.
14 List of Figures
Figure 1. Temperature dependence of equilibrium amount of all phases
in one mole of Alloy 825 “FCC” is a nickel-rich face-centred cubic phase, “BCC” is a body-centred cubic phase with the approximate
composition M(C,N)3, “sigma” and “laves” are intermetallic phases,
other phases have the approximate chemistries specified in the corresponding labels. ... 32
Figure 2. Flow chart for how all supplements are related to each other.
... 34 Figure 3. Micro-etched sample of as-cast Alloy 825 starting material used for hot compression test at a magnification of 200x a) Columnar grain structure-transverse b) Equiaxed grain structure-transverse. ... 45
Figure 4. Schematic diagram of thermal profiles used in the
thermomechanical tests (applied to Supplement II). ... 46
Figure 5. Work hardening rate (solid line) and true stress (dash line)
as a function of true strain for temperature of 1200 ℃ and strain rate
of 1 s for Alloy 825 with an addition of magnesium (supplement
II). ... 49
Figure 6. Schematic diagram showing thermo-mechanical processing
cycle and soft annealing. “min mm-1” refers to the heat treatment time
per millimetre of rod radius. ... 51
Figure 7. EBSD images of (a) and (b) continuously cast bloom
specimen of Alloy 825 without magnesium; (c) ingot cast specimen of Alloy 825 with magnesium; (d) wrought specimen parallel to forging direction (FD) after solution annealing at 1200 °C. ... 67
Figure 8: Optical microstructures at the centre area of the columnar
(C) and equiaxed (E) samples after completion of the hot deformation tests at the indicated temperatures, to a total nominal strain 0.7 and at a strain rate of 1 s . ... 69
Figure 9: EBSD map of Alloy 825 after a uniaxial compression to
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deformation temperature. Blue regions represent fully recrystallized grains, Yellow regions refer to grains that have undergone recovery, red regions represent areas that have undergone neither recovery nor recrystallization. The images are overlaid with high angle grain boundaries in black and low angle grain boundaries in white. ... 69
Figure 10: (a) Variation of the percentage of recrystallized grains and
recrystallized grain sizes in the columnar and equiaxed grains structures (b) hardness (Hv) of deformed columnar/equiaxed Alloy 825, with error bars representing the standard deviation. Data are from the deformation temperature over a range of 800 °C to 1200 °C and at nominal strain of 0.7 for an Alloy 825 after a hot compression using a strain rate of 1 s-1 and a strain of 0.7. ... 71
Figure 11: True stress-true strain curves for Alloy 825 tested at strain
rates ranging from 0.1 s-1 to 10 s-1 and temperatures of (a) 1150 °C and (b) 1200 °C. The vertical order of the dashed lines is the same as for the solid lines. Solid lines are for samples that contain magnesium; dashed lines represent data for samples to which magnesium has not been added. To aid a clear comparison, the -axis range is kept constant in both subfigures. ... 73
Figure 12. Micrographs after the hot deformation at 0.1 s−1 of the
materials and at temperatures of 1100 °C (a) Alloy A (b) Alloy B. The fraction of recrystallized grains in Alloy A is higher than that of Alloy B at all deformation temperatures. ... 75
Figure 13. EBSD orientation images of the microstructure for Alloy A
and B deformed at (a) Alloy 825 with the addition of magnesium deformed at 1150 °C at a strain rate of 0.1 s−1 and (b) conventional
alloy 825 deformed at 1200 °C with a strain rate of 1.0 s−1. Each image
is overlaid with high angle grain boundaries (white) and low angle grain boundaries (black). CD means the direction parallel to the direction of casting the length of the continuously cast strand, or the long direction of the ingot, as appropriate. ND means the orientation normal to the casting direction, CD. ... 75
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Figure 14. Effect of deformation temperature on the fraction of
dynamically recrystallized grains. The error bars represent the standard error about the mean value. ... 76
Figure 15. Microstructural evolution following hot-deformation at a
strain rate of 0.1 s and at temperatures of (a) and (b) 1100 ℃, and at
a strain rate of 10 s (c) and (d) 1250 ℃. Each image is overlaid with
high angle grain boundaries (white) and low angle grain boundaries (black). All images are presented using the same scale. The colour map of the pole figure maps shown in (a)–(d). The regions displayed are representative of the entire material. CD means the direction parallel to the direction of casting the length of the continuously cast strand, or the long direction of the ingot, as appropriate). ND means the orientation normal to the casting direction, CD. ... 79
Figure 16. Relation between (a) the peak strain and (b) dynamically
recrystallized grain size with the Zener–Hollomon parameter, . The standard error of the mean dynamically recrystallized grain size in measurements was calculated and presented as the error bars. ... 80
Figure 17. Variation of peak stress, ln P, versus strain rate, ln at
different deformation temperatures. This plot is used to derive the materials parameters and . ... 84
Figure 18. Power dissipation maps for as-cast Alloy 825 at a true strain
of 0.7. The contour values show the efficiency of power dissipation in percent. ... 85
Figure 19. EBSD image (band contrast) of equiaxed compressed
specimen of an Alloy 825 after a uniaxial compression to a nominal
strain of 0.7 and a strain rate of 10 s at the deformation temperature
of 950°C, each overlaid with high-angle (black) and low-angle (white) grain boundaries. ... 86
Figure 20. Hardness map of hardness with strain rate and temperature
at 0.7 strain. ... 87
Figure 21. The effect of strain magnitude on the fraction of refined
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Figure 22. Effect of the hot forging strain on the recrystallized
austenite grain size (open triangles) and calculated (dashed line) values in Alloy 825 samples (supplement IV). ... 91
Figure 23. The effect of forging strain level on the dislocation densities
of experimental (solid triangle) and calculated (line, Equation 24) values in the present Alloy 825. ... 92
Figure 24. Contribution of different strengthening mechanisms to
general yield strength of hot forged Alloy 825 subjected to different strain levels. ... 94
18 List of Tables
Table 1. Composition in wt.% for the specification for Alloy 825
(ASTM B425). ... 30
Table 2. Requirements for Ni-Fe-Cr-Mo-Cu Alloy (UNS N08825)
according to ASTM B425-11 in the form of annealed hot-finished rounds ... 31
Table 3. Overview of the supplements I-V ... 35
Table 4. Chemical compositions of alloy 825 in this study. All
compositions expressed in wt.%. Levels of carbon and sulfur were measured using combustion photometry. X-ray fluorescence spectrometry was used for other elements. Uncertainty Uncertainty estimates for C and N measurements are taken from data in standard ASTM E1018-11. The uncertainty estimates of the other elements are taken from ASTM E572-13. ... 44
Table 5. Deformation temperature, nominal strain rate, and nominal
true strain used in this thesis. ... 48
Table 6. Sample designations used in the present work. ... 51
Table 7. A scanning of step size, covering area and grain orientation
spread ... 54
Table 8. Fitted parameters and for the peak strain and stress in both
alloys where the fit is of the form !. ... 80 Table 9. Materials constants used in Equation 15, derived from regression analysis, including the standard error of regression. ... 84
Table 10: Extent of recrystallization in selected samples from the
current study. Both high temperature and low strain rate promote recrystallization. ... 86
Table 11. Strengthening coefficients for solid solution strengthening
in nickel [70,75] ... 96
Table 12. Calculated Orowan bowing strengthening contributions
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Table 13. Overview of the results and applications of the supplements.
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List of Symbols
Supplements I-II
Zener–Hollomon parameter, s-1
"D Mean dynamically recrystallized grain size,
µm
% Material constant
& An exponent in the power law function for
Zener-Holloman parameter,
σp Peak stress, MPa
εp Peak strain
Strain rate, s-1
* Work hardening rate, MPa, (θ =-.
-/01,/)
M Standard error of the mean value, %
* The minimum misorientation angle,
degrees, EBSD
"2 Initial grain size, µm
and Fitted parameters for the peak strain and
stress
3 and & Dimensionless material constants used with
an Arrhenius-type equation
&′, 5 and 6 Adjustable parameters, MPa-1
7p, 7E and 7H Activation energy for deformation assuming
a power, an exponential and hyperbolic-sine equation
8 Universal gas constant, 8.31 J mol-1 K-1
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: Power law exponent, dimensionless
constant
; Units of length, μm
Supplement III
True strain
< Strain rate sensitivity, m
= A total power applied to the material during
deformation
> Some power consumed to perform the
deformation and cause any temperature change
? Remaining power stored as changes to the
microstructure
@ Efficiency of power dissipation, %
A, and c Temperature-dependent strength
parameters
B Metallurgical instability parameter during
plastic flow Supplement
IV
8C Reduction Ratio
KAM Kernel Average Misorientation, degrees
*KAM Kernel Average Misorientation Angle,
degrees
D Dislocation density, m-2
D2 Initial dislocation density at solution
anneal, m-2
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F Step size of the electron backscatter
diffraction scan, µm
G Scanning geometry constant
HSFE Stacking fault energy, mJ m-2
Y Yield Strength, MPa
UTS Ultimate Tensile Strength, MPa
f Strain to Failure; Elongation, %
Σ3 Twin boundaries
Fraction of recrystallized grains, %
K and & Material constants and depend on the grain
size
LV Vickers hardness, MPa
K Grain boundary strengthening factor, MPa
m0.5 (Petch-coefficient)
0 Inherent resistance of the material to
dislocation glide
5 or 5M A proportionality constant or substructural
strengthening factor
G Shear modulus of the material, 76000 MPa
M Taylor factor, (FCC~3.1)
6 and & Materials constants
Supplement V
Np and Op The volume fraction, % and average
precipitate radius, nm
p Peierls stress of the matrix phase, MPa
s Solid solution strengthening contribution,
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O Orowan bowing around precipitates in the
matrix, MPa
g Grain refinement, MPa
w Work hardening, MPa
PQ Solution strengthening coefficient of the Rth
solute
SQ The atomic fraction of the Rth solute
>Ni Shear modulus of nickel, MPa
TQ The lattice strain
U2 Critical resolved shear stress for dislocation
motion, MPa
V A constant characterizes the interaction
between dislocations and the precipitates
W Mean spacing between precipitates, µm
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List of Abbreviations
AOD Argon Oxygen Decarburization
TEM Transmission Electron Microscopy
STEM Scanning Transmission Electron Microscope
SEM Scanning Electron Microscopy
EDS Energy Dispersive X-ray Spectroscopy
EBSD Electron Backscattered Diffraction
XRD X-ray Diffraction
GOS Grain Orientation Spread
CSL Coincidence Site Lattice
DRX Dynamic Recrystallization
SRX Static Recrystallization
HAGB High-angle grain boundary
LAGB Low-angle grain boundary
ND Normal direction in a compressed sample
RD Rolling/radial direction in a compressed sample
TD Transverse/tangential direction in a compressed
sample
CD Compression direction in a compressed sample
FD Forging Direction
PSN Particle-stimulated nucleation of
recrystallization
OIM Orientation Imaging Microscopy
TSL Tex Sem Ltd.
FCC Face Centred Cubic phase
BCC Body-Centred Cubic phase
ASTM American Society for Testing and Materials
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PHS Precipitation hardened superalloy
SS Solid solution
CW Cold worked nickel-base alloy
26 Contents
1. Introduction ... 28 1.1. Objectives of the work ... 33 1.2. Present work ... 37
1.3. Contribution to sustainable development ... 40
2. Experimental and analytical methods ... 43
2.1. Material ... 43 2.2. Sample preparation for thermo-mechanical tests ... 44 2.3. Thermo-Mechanical Testing ... 45 2.4. Peak stress/strain estimation ... 48 2.5. Strain magnitude during hot forging process... 49 2.6. Optical microscopy sample preparation ... 51 2.7. EBSD sample preparation... 52 2.8. Electron Backscatter Diffraction Analysis (EBSD) .... 52 2.9. Tensile specimens and testing ... 55 2.10. Hardness testing ... 55 2.11. Transmission electron microscopy ... 56 3. Modelling ... 57
3.1. Modelling of flow stress and microstructural behaviour
57
Flow stress model (Determination of material constants) ... 57 Dynamically recrystallized grain size model ... 59 Processing map (Dynamic Material Model) ... 59 3.2. Modelling forging microstructure and internal stresses
61
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Structural strengthening ... 62 3.3. Modelling of strengthening mechanisms in Alloy 825
63
4. Results and discussion ... 67
4.1. Microstructure characterization in cast
equiaxed/columnar Alloy 825 ... 68 4.2. Structure and property comparison of Alloy 825 with/without Mg ... 72 Hot deformation stress behaviour ... 72 Microstructural evolution ... 74 Crystallographic texture... 77 Peak strain/stress and microstructural behaviour ... 80 4.3. Processing maps of cast Alloy 825 ... 83 Establishment of processing map for Alloy 825 ... 83 Characterization of microstructure in different domains of the processing map ... 87 4.4. Forging microstructure and internal stresses ... 89 4.5. Structural strengthening ... 92 4.6. Solid solution and Orwan bowing (precipitates) strengthening ... 95 5. Concluding discussion ... 98 6. Conclusions ... 103 References ... 107
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1.
IntroductionIn the first part of the thesis a literature review is presented to provide the reader with basis for further reading of this thesis. Alloy 825, the specific alloy being the focus at this study, with its excellent combination of mechanical properties and corrosion resistance [1–3], can be used at high temperatures and in acidic environments. It is used in the petrochemical sector, such as in the oil and gas extraction and petroleum refining industries. It is also used in other applications to make tanks that are subject to corrosive environments. In these applications, the alloy is subjected to both mechanical loading and prolonged contact with corrosive substances. Therefore, both the mechanical properties and corrosion resistance can be limiting. If both the mechanical properties and corrosion resistance can be improved, the service life of components made from the alloy can be increased. This will promote a sustainable development by reducing the resource consumption and increasing the economic benefits of using the component.
Alloy 825 is utilized both for wrought cold worked and hot finished annealed conditions. The wrought forms of Alloy 825 products are either continuously cast or ingot cast. The continuously cast products often contain large columnar grains adjacent to the mould wall and fine equiaxed grains in the centre. However, a uniform microstructure is desirable to maximise the mechanical and corrosive performance in all directions and in all regions of a component. The casting structure is broken down by thermo-mechanical processing, either forging or rolling at high temperature to obtain a uniform chemistry and microstructure. Previous work recommends that hot working should be performed at deformation temperatures
between 870 °C and 1180 °C to achieve a dynamic
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Deformation is commonly by hot forging, but there are few data about how hot forging affects the microstructure, which is critical for both the mechanical and corrosion performance. Industrial forged products, particularly billets or bars, manufactured from the alloy are subjected to high deformation temperatures and are exposed to deformation schedules, which are varied to meet property requirements. However, owing to manufacturing difficulties, the microstructure in as-delivered materials varies widely, due to differences in the grain size. Hence, the mechanical properties in different heat lots have been shown to be different. These differences are caused by the differences in the grain structure, especially (whether) come to if the samples are taken from the columnar or equiaxed regions of the cast ingot/strand, and/or by hot-working conditions such as the deformation temperature, strain and strain rate. Thus, the microstructure of the alloy 825 during hot deformation can be controlled by controlling the thermomechanical process and by promoting a dynamic recrystallization (DRX). It is well known that the addition of microalloying elements such as magnesium, calcium and boron have a strong influence on the mechanical properties of both wrought and cast nickel-base superalloys. Specifically, they can promote a dynamic recrystallization and widen the process window that will result in the desired microstructure [4–6].
The high contents of nickel, chromium and molybdenum in alloy 825 give a good corrosion resistance and improves the mechanical properties such as the yield strength, ultimate tensile strength, and elongation to failure [1,2,7–9]. According to industrial standards, the target for the hot-annealed material
is an alloy with a yield strength, y ≥ 241 MPa [10]. The alloy,
which also contains small additions of titanium, is expected to form a matrix of a face-centred cubic (austenitic) phase, containing small amounts of titanium carbonitride. After being
30
deformed, the material is solution annealed (or stabilized) at a suitable temperature to precipitate the maximum volume fraction of Ti(C, N) inside grains and to avoid the precipitation
of Cr23C6-type carbides at grain boundaries at later stages of the
thermomechanical treatment [11,12]. Cr23C6-type carbides
would deplete the matrix of chromium, which could lead to a sensitivity to corrosive attack [11,13]. Mechanical properties in any alloy, are affected by solid solution strengthening, which is provided in alloy 825 by the presence of iron, chromium and molybdenum, strain hardening, and the grain size [14]. For this reason, the alloy should be carefully processed to produce a small grain size and a maximize strength using grain refinement. It is well known that thermo-mechanical processing in combination with alloying, precipitates, or initial grain size in any alloy have an influence on the dynamic recrystallization during hot deformation [13,15]. For annealed hot-finished rounds, the minimum mechanical requirements of the standard specification for Ni-Fe-Cr-Mo-Cu Alloy (UNS N08825) Rod and Bar, ASTM B425-11 are outlined in Table 1 and Table 2 [23]. Table 1. Composition in wt.% for the specification for Alloy 825 (ASTM B425).
C Si Mn Cr Fe Mo Ti Cu Ni
Min. 19.5 22.0 2.5 0.6 1.5 38.0
31
Table 2. Requirements for Ni-Fe-Cr-Mo-Cu Alloy (UNS N08825) according to ASTM B425-11 in the form of annealed hot-finished rounds Condition Tensile Strength, Min. [MPa] Yield Strength 0.2 % offset, Min. [MPa] Elongation Min. % Annealed: Hot-finished 586 241 30
Thermodynamic modelling on Alloy 825 using Thermo-Calc modelling software with the TTNI8 database has been published [16]. These results were further supported through calculations of the equilibrium phase fractions as a function of temperature performed based on the measured composition (Table 1, Supplement V) using Thermo-Calc 2020a and the TCNI9 thermodynamic database. The alloy is expected to form a matrix of face-centred cubic phase, containing small amounts of titanium carbonitride and alumina in the temperature range 900 < 9 / ℃ < 1300 (Figure 1) [11]. Furthermore, the harmful phase existing in this alloy was clearly seen as being the σ phase (“Sigma”). This phase has a tetragonal structure which does not fit well with the matrix. It forms after a prolonged time at temperatures between 540 and 980 °C [17,18]. The sigma phase is expected to be thermodynamically stable between 600 °C and 925 °C. Above this range, it is unstable, and any precipitates will dissolve. Also, a small amount of TiN phase can precipitate below 1550 °C. According to the phase diagram [16], an alloy 825 would be stabilized against intergranular attack if the material is annealed at 940°C [1]. Also, the small titanium carbonitrides are the only phase that will be considered to cause coarsening in this
study: the face-centred cubic is the matrix phase. A previous
study of expected phases in Alloy 825 (referred to by the alternative name UNS N08825) on a standard composition
32
using thermodynamic modelling were almost identical to the current calculation results (Figure 1), although no alumina phase was reported [16]. There is no apparent explanation for this, but it is possible that some phases such as alumina were disabled in earlier simulations. Given the very small volume fraction expected, it will be assumed that alumina does not form in significant amounts. Therefore, it may be ignored unless proven otherwise.
'
Figure 1. Temperature dependence of equilibrium amount of all phases in one mole of Alloy 825. “FCC” is a nickel-rich face-centred cubic phase, “BCC” is a body-face-centred cubic phase with
the approximate composition M(C,N)3, “sigma” and “laves” are
intermetallic phases, other phases have the approximate chemistries specified in the corresponding labels.
33
1.1. Objectives of the work
The aim of the present work is to identify the microstructure of alloy 825 after single-hit hot compression experiments. The work also investigates the influence of different total strains on the microstructure and mechanical properties to have direct relevance to the current industrial production processes. In Figure 2 a schematic overview of how the supplements are connected is presented.
The experimental part of this work has been performed by using
laboratory-based thermomechanical testing. Thereafter,
observations and evaluations of the microstructure following a heat treatment were carried out by using optical microscopy and electron backscatter diffraction techniques. Observations and evaluations of the microstructure and resulting mechanical properties were carried out using tensile testing, optical microscopy, SEM and EBSD. In Table 3 an overview of the objectives, approaches and parameters used in the different supplements is presented.
34
Figure 2. Flow chart for how all supplements are related to each other.
35
Table 3. Overview of the supplements I-V
Suppl Objective Approach Parameters
I Study the microstructural evolution of cast alloy 825 due to dynamic recrystallization
Laboratory analysis of the true stress and Vickers hardness;
characterization of samples deformed at different temperatures using optical microscopy and EBSD Thermomechani cal testing, Data from optical microscopy and EBSD
II Study the effect of trace additions of magnesium on the mechanical properties of cast alloy 825 with emphasis on whether or not dynamic recrystallization occurs
Optical microscopy and SEM/EBSD analysis of microstructure following compression at different deformation temperatures Tensile tests and analysis of the peak true stress-strain used to model the Zener-Hollomon
parameter, Z, and indicate the onset of dynamic recrystallization Thermomechani cal testing, tensile testing, optical microscopy and EBSD III Derive processing map for hot deformation of alloy 825; relate microstructure to mechanical properties Dynamic material modelling used to model the processing conditions for Ni-based alloy 825, based on experimental data Thermomechani cal testing, optical microscopy, EBSD, Vickers hardness; derived process map for alloy 825
IV Relate
microstructural changes to total
Industrial trials of plastic deformation at elevated temperature designed to
Industrial trial, TEM, SEM, EBSD, tensile
36 deformation strain stimulate dynamic recrystallization testing, Vickers hardness V Quantify the strengthening mechanisms in Alloy 825 during soft annealing and compare with calculated values to tensile test data
Compare results of tensile tests to microstructure and theoretical strength contributions to quantify each contribution to the overall strength in the alloy Literature review, data from industrial trial, SEM, EBSD, tensile testing, Vickers hardness, thermodynamic modelling
37
1.2. Present work
This thesis consists of 5 supplements and the detailed description of each supplement is given as follows:
Supplement I reports the results from the optical and EBSD images of both columnar and equiaxed samples observed after deformation of alloy 825 at various temperatures after single-hit compression tests. The true stress-true strain curves of both structures recorded by thermo-mechanical testing are discussed. The effect of different testing conditions and various deformation temperatures on the microstructure and micro-hardness in both structures are discussed. Very little dynamic recrystallization was observed in both columnar and equiaxed structures for all tested deformation temperatures. Instead, a dislocation substructure consisting of low-angle grain boundaries was found to dominate.
Supplement II studies the effect of trace additions of magnesium on the mechanical behaviour of cast alloy 825 after one-hit deformation, with emphasis on whether or not a dynamic recrystallization occurs. The true stress-true strain curves were recorded for cast alloy 825 with and without the addition of magnesium, following single-hit compression tests. In addition, the effect of different testing conditions and various deformation temperatures on the microstructure and microhardness in both structures are discussed. The results from the optical and EBSD images in both alloys structures for the various deformation conditions were presented. The structural changes in both alloys were characterized by the development of continuous and discontinuous dynamic recrystallization for any test conditions. If magnesium is added, the fraction of the high angle grain
38
between 43% and 70% higher in alloy 825 for all the temperatures and the strain rates being tested.
Supplement III reports the studying of the hot-working behaviour of cast Alloy 825 for a wide range of processing conditions. The final deformed microstructure was compared to the mechanical properties during the hot-forming process to develop a processing map, based on dynamic materials modelling. This allowed the deformation of an optimum hot-working processing windows for Alloy 825 and the microstructure in the hot-worked material to be optimised. The true stress-true strain curves of cast Alloy 825, measured during hot compression, are discussed. The effect of different testing conditions, various deformation temperatures and strain rates on the microstructure and microhardness in both structures are discussed. The results showed that the flow stress is both depending on the strain rate and temperature. Furthermore, It increases with an increased strain rate and a decreased temperature. The peak true flow stress of Alloy 825 can be obeyed by the Zener-Hollomon equation. In the process map, a
domain at a deformation temperature of between 1100 ℃ and
1250 ℃, with a strain rate of between 0.01 s and 0.1 s gives the optimum hot deformation condition. It exhibits a high-power dissipation efficiency (up to 35%) and can result in the formation of a fine recrystallized microstructure. The hot processing map suggests a good workability for a wide range of
conditions with a strain rate, < 0.1 s and a temperature of
1100 ℃ or higher. The optimum processing parameters to reach a good strain hardening value are obtained in the temperature
range between 950 ℃ and 1100 ℃ and using a strain rate of
39
Supplement IV reports the microstructural evolution and their influence on the yield strength of Alloy 825 processed by different strain forging levels. SEM, EBSD and TEM images are presented, analysed, and compared to tensile test data. The influence of strain magnitude on the average grain size, microhardness, and dislocation density in Alloy 825 deformed samples. The results revealed that the average grain size decreases with an increasing strain during forging, due to an increased recrystallization. Both continuous and discontinuous dynamic recrystallization mechanisms operate during the hot
forging process. The area fraction of recrystallized grains, G,
with sizes below 25 µm, which have a GOS threshold value ≤ 1°,
increased with an increasing strain, . The fraction of grains that are recrystallized can be described using a simplified modification of the Johnson–Mehl–Avrami–Kolmogorov equation:
G = 1 − exp_−1.265 2.`ab.
The experimental 0.2% proof strength, 2.c, may be obtained by
multiplying the calculated yield strength by a factor of 1.34 and
can also be expressed by using the initial grain size, d2, and the
total forging strain, , by using modified Hall-Petch relationship:
2.c,efge= 193 + id2exp_−2 bm 2.no0.42
+ 1.3 × 10 `id
2exp_−2 bm p
2.c,qrs = 1.34 2.c,efge
The maximum yield strength and ultimate tensile strength were obtained after forging to a true strain of 0.9 and were 413 MPa and 622 MPa, respectively, with a ductility of 40%.
Supplement V quantifies the strengthening mechanisms in a commercially produced sample of Alloy 825 during soft annealing (stabilization) and compares these findings to tensile
40
test data. The study investigates the effect of the forging strain magnitude and subsequent annealing on the microstructure, strengthening mechanisms, and room temperature mechanical properties to assess the suitability of the current industrial practice. The microstructural changes and mechanical properties were analysed using SEM-EBSD and tensile testing. Thermokinetic simulations were also used to calculate the interfacial energy between the precipitate and the matrix for modelling of strengthening mechanisms. The yield strength, ultimate strength and ductility are 348 MPa, 618 MPa, and 40% respectively. Also, solid solution strengthening is very significant in the current alloy, contributing 127 MPa to the yield strength. Precipitation strengthening in the current alloy was found to be 48 MPa and it can be predicted to a good accuracy (46 MPa, 4% difference) using the conventional Orowan equation. The grain refinement and work hardning strengthening was evolved by
forging strain levels and they were found to be between 26 MPa
and 35 MPa and 25 MPa and 43 MPa respectively. The
interfacial energy of Ti(C,N) in the face-centred cubic matrix of
the current alloy is 0.8 mJ m c. It is derived by simulating
precipitate and growth of Ti(C,N) during the soft annealing treatment.
1.3. Contribution to sustainable development
The work aims to realize breakthrough improvements in resource efficiency in the manufacturing of Alloy 825 long products, which are key engineering material in many types of demanding applications. Today, a low tensile strength is a major source of yield loss and process inefficiency in producing Alloy 825 long products. Optimised processes that minimize negative environmental impacts, conserve energy and natural resources are needed. This also enhances employee, community, and
41
product safety (SDG 8 and SDG 9). Based on this work, provided in the five supplements potential yield improvements and production cost reductions are identified, if a simplified production route producing high tensile strength products could be applied to a commercially important Alloy 825 (SDG 8). Improved mechanical properties of Alloy 825 material will enable a more efficient manufacturing process, which reduces energy efficiency input and therefore the CO2 emissions for manufacturing reduces (reduced carbon emissions) (SDG 13). It also reduces the chance that a component will fail, which reduces environmental impact even more (SDG 13). It also makes the manufacturing process more profitable, by reducing costs for energy and failed manufacturing runs (SDG 8).
The production of Alloy 825 and other alloys is mainly based on recycled steel scrap. With the optimized use of secondary raw materials, the need for costly primary raw materials is reduced. By using recycled material in our production processes, we can increase our contribution to a more circular society (SDG 9). Alloy 825 material is attractive in a market due to that it is widely used in most industrial settings, e.g. tanks, vessels, agitators, downhole equipment, and nuclear reprocessing where aggressive chemical environments are present and where the use of other alloys have failed. There is a growing demand for chemical-based products and processes, which are major contributors to national and world economies. Component failures raise costs and can lead to leaks and leakages of chemicals, which harms both health and the environment. Product life cycles are essential in order to avoid significant and increasingly complex risks to human health and substantial costs to national economies. The implementation of the results from this thesis will enable component life to be extended by
42
optimising the microstructure, corrosion resistance and mechanical properties (SDG 12).
Heating Alloy 825 material to higher temperatures requires more energy to heat the material, but it is not necessarily bad for a sustainable development. This is due to that it allows more a greater fraction of mechanical energy to be used constructively during the production of Alloy 825, according to the current findings (SDG 7 and SDG 12). The electricity used to perform this heating is obtained from clean sources to maximise the sustainability.
43
2.
Experimental and analytical methodsA Gleeble 3500 thermo-mechanical simulator was employed to investigate the influence of the deformation temperature and different initial grain structure on the flow stress. The dynamic recrystallization behaviour in hot deformed cast alloy 825 samples, both in the columnar region adjacent to the mould wall and in the equiaxed region in the centre to the mould wall were
studied by using EBSD. Tensile testing using a strain rate of 10-3
s-1 was performed in the R&D centre at AB Sandvik Material
Technology in Sandviken, Sweden. Also, optical microscopy and SEM investigations were performed at the Sandvik Materials Technology R&D centre.
2.1. Material
All material in this study has been provided by Sandvik Materials Technology. The materials that were used for the investigation in this work are listed in Table 4. Samples without the addition of magnesium (Supplement I and II, Table 4) were produced by AOD-refining, followed by continuous casting. Samples that included magnesium additions (Supplement II, Table 4) were manufactured by argon oxygen decarburization (AOD) refining and followed by ingot casting.
44
Table 4. Chemical compositions of alloy 825 in this study. All compositions expressed in wt.%. Levels of carbon and sulfur were measured using combustion photometry. X-ray fluorescence spectrometry was used for other elements. Uncertainty Uncertainty estimates for C and N measurements are taken from data in standard ASTM E1018-11. The uncertainty estimates of the other elements are taken from ASTM E572-13. Structure/ Supplement C Cr Mo Ti Cu N Mg Fe O Cont. Cast/ I, III 0.007 22.0 2.53 0.80 1.60 0.01 0 32 7 ppm Ingot Cast/ II 0.020 22.2 3.09 0.76 1.58 0.012 0.0076 30 10 ppm Wrought/ IV, V 0.02 22.0 3.0 0.70 1.8 0.018 0 30 10 ppm Uncertainty 0.01 0.001 0.03 0.003 0.002 0.005 0.0001 - -
2.2. Sample preparation for thermo-mechanical tests
The cylindrical compression samples, 15 mm high and 10 mm diameter, were cut and machined with the compression axis parallel to the long axis of columnar grains and in the corresponding direction in equiaxed regions of the cast. Figure 3 shows the single-phase austenitic microstructure of the as-cast Alloy 825 starting material, etched using a Beraha 9b solution. The columnar and equiaxed samples were not homogenized and they were etched using a mixture of 85 ml water, 15 ml
hydrochloric acid and 0.6-1 g potassium metabisulfite (KcScOn).
The ‘columnar’ structure characterises as elongated grains which consists of a fine network of thick dendrite arms and a small number of grain boundaries [19]. Conversely, the ‘equiaxed’ structure is the very fine central zone, which consists of a large number of thin dendrite arms and a large number of fine grains. Hence, it contains a greater density of grain boundaries and a small amount of porosity.
45
Figure 3. Micro-etched sample of as-cast Alloy 825 starting material used for hot compression test at a magnification of 200x a) Columnar grain structure-transverse b) Equiaxed grain structure-transverse.
Samples for microstructural analysis taken from the columnar region exhibited a random orientation, which contain small TiN inclusions. Furthermore, those taken from equiaxed grains contain a small amount of porosity.
2.3. Thermo-Mechanical Testing
A Gleeble 3500 test machine and a compatible equipment, equipped with a computer, were used to capture stress-strain data during hot compression tests [20]. Cylindrical compression
test pieces with a diameter of 10 mm and a length of 15 mm in
cast Alloy 825 without Mg (supplement I and II) were machined with the compression axis parallel to the long axis of the continuously cast blooms and in a region where grains were found to be columnar. Table 5 shows the designations of all the experiments for supplement I, II and III. Deformation tests were carried out at a range of deformation temperatures from 800 °C to 1200 °C with a total nominal strain of 0.7 and a
constant strain rate of 1 s , (Supplement I, Table 5). Hot
compression test pieces of Alloy 825 alloying containing Mg (supplement II) were produced with the same dimensions and
46
the compression axis parallel to the ingot axis from a constant radial position. Furthermore, where columnar grains occurred. Tests were carried out under two conditions: a constant
temperature in the range of 1200 ℃ and 1250 ℃, and at two
different temperatures of 1100 ℃ and 1150 ℃ (Figure 4).
Figure 4. Schematic diagram of thermal profiles used in the thermomechanical tests (applied to Supplement II).
These conditions were selected as the dynamic recrystallization was not complete and the flow stress was not affected by the cast grain structure or grain size [21,22]. Figure 4 shows a schematic diagram of the thermo-mechanical tests. Samples tested at two temperatures were initially heated to the homogenization
temperature of 1200 ℃ for 100 s. Therefore, the samples were
cooled to the deformation temperature at a rate of 5 ℃ s . The
samples were held at each deformation temperature for 30 s.
47
vacuum, before being quenched in high pressure air to a
temperature of between 100 and 200 ℃.
To derive the processing maps of alloy 825, hot compression tests were conducted to a total nominal strain 0.7 at strain rates of between 0.01 s and 10 s and with intervals of an order of magnitude as well as using deformation temperatures between 950 °C and 1250 °C with intervals of 50 °C (Supplement III). In
the case of deformation at 1200 ℃ or 1250 ℃, samples were
soaked at the deformation temperature for 100 s and then
deformed. In the case of samples deformed at temperatures
below 1200 ℃, samples were first soaked at 1200 ℃ for 100 s to
mimic an industrial practice. Therefore, they were cooled at 5 ℃ s to the deformation temperature. These conditions were selected as it has been shown that the dynamic recrystallization is unlikely to be complete and the flow stress is unlikely to be affected by the cast grain structure or the grain size [21,22]. In supplements I, II, and III, the samples were held at the
deformation temperature for 30 s before deformation to ensure
that a homogeneous temperature were reached to ensure all precipitates were dissolved (except titanium nitride). Samples were subsequently cooled to the deformation temperature at 5
°C s , and held at each deformation temperature for
approximately 30 s to ensure that a uniform temperature had
been reached. Samples were compressed by a load of 100 kN
under vacuum, before being quenched in high pressure air to a
temperature of between 100 ℃ and 200 ℃. Table 5 shows the
48
Table 5. Deformation temperature, nominal strain rate, and nominal true strain used in this thesis.
Experiment Deformation Temperature, [°C] Nominal Strain Rate, [s-1] Nominal True Strain Supplement I 800, 900, 1000, 1100,1200 1.0 0.7 Supplement II 1100, 1150, 1200, 1250 0.1, 1.0, 10.0 0.7 Supplement III 950, 1000, 1050, 1100, 1150, 1200, 1250 0.01, 0.1, 1.0, 10.0 0.7
2.4. Peak stress/strain estimation
The measurement of peak stress and strain was achieved using
the measured work hardening rate, * =xy
xM0z,M [23], which was
calculated from the stress and strain measurements by the averaging the slopes of three adjacent points for each data point (Equation 1). To overcome noise in the data, the stress–strain curves were fitted using a high-order polynomial [24–26].
* = 0.5_ {| {
{| − {+
{ − {
{− { b Equation 1
During the strain test, the strain hardening coefficient, * ,
decreases quickly as the stress value, , increases. This means that the rate of increase in the flow stress decreases. Eventually,
the lines of * and cross. The conditions at this time are defined
as the critical stress, c, and critical strain, c. The position at
which the lines cross depends on the axis scales of each quantity. Therefore, the critical stress and strain can change depending on the way in which the graph is plotted. However, these points are
49
not used for subsequent analysis, so this issue is not problematic in this study. During a subsequent deformation, continues to
increase and * is reduced to zero. The stress at which this occurs
is defined as the peak stress, p. Similarly, the corresponding
strain is defined as the peak strain, p. * then decreases to a
minimum value before rising again. When it increases to zero,
the stress is called the steady state stress, } with a
corresponding steady state strain, s. Both the peak and steady
state properties depend on the absolute values of the measured quantities and are therefore insensitive to axis scaling. The definitions of the points are summarized in Figure 2 [21].
Figure 5. Work hardening rate (solid line) and true stress (dash
line) as a function of true strain for temperature of 1200 ℃ and
strain rate of 1 s for Alloy 825 with an addition of magnesium
(supplement II).
2.5. Strain magnitude during hot forging process
Wrought material of Alloy 825 was cast after air melting in an electric arc furnace and refined in an argon oxygen decarburization process (Supplement IV and V,
50
Table 4). All material in Supplement IV and V came from three wrought billets of Alloy 825, which originated from the same cast ingot. The ingot was then subjected to hot-forging at 1200 °C with a total 70% reduction to obtain three billets, which served as the starting billets for subsequent experiments. The billets were hot forged at 1200 °C to true strains of 0.45, 0.65 or 0.90, (Figure 6, Table 6). The billets were soaked at 1200 °C for a time equal to 3 min per millimetre of billet radius and then hot forged. The forging process was performed using a hydraulic press with
flat dies and using a strain rate of ~0.5 s . The hot-forging of all
three billets was performed at temperatures maintained
between 950 ℃ and 1180 ℃ (Figure 6). Samples were reheated
during each forging process and the final forged bars were
quenched in water from between 950 ℃ and 980 ℃. Different
samples were subjected to total accumulative strains of 0.45,
0.65 or 0.9 with a pass strain of ~0.1 (i.e. 10% reduction per pass) to study the structural changes during deformation (Table 6). The samples were rotated by 90° from one pass to the next.
The true strain was estimated by using the formula ≈ ln 8C,
where 8C is the reduction ratio (ratio of the starting
cross-sectional area to the final cross-cross-sectional area).
The alloy production and processing took place at Sandvik Materials Technology facilities in Sandviken, Sweden. The material was sectioned for microscopy parallel to the forging (axial) direction from the centre of the solid bar. After forging, the bars were solution annealed at a temperature of 950 °C for 1.5 min per millimetre of bar radius. From each annealed bar, samples were extracted for microscopy in the parallel to the forging (axial) direction from the centre of the solid bar.
51
Table 6. Sample designations used in the present work. Sample designation Solution annealed A B C True strain 0.00 0.45 0.65 0.90
Figure 6. Schematic diagram showing thermo-mechanical
processing cycle and soft annealing. “min mm-1” refers to the
heat treatment time per millimetre of rod radius.
2.6. Optical microscopy sample preparation
The deformed specimens were sectioned through the longitudinal axis and prepared for optical microscopy using standard metallographic techniques with a final stage of using a diamond polishing to 3 µm. The sectioned surfaces were electrolytically etched in a solution of 10 g oxalic acid and 100 ml water and using a voltage of 6 V for 3-60 s. Micrographs were
52
taken from the centre of the longitudinal sections. Samples were also used for hardness tests, so the non-recrystallized grain structure could be observed together with any recrystallized grains. The samples were tested for hardness on a Vickers hardness tester, using a 1.0 kg load. Each sample was tested 10 times. The reported measurement values are the mean of those 10 measurements with the corresponding standard deviations. All measurements were taken on the same equipment and followed the same procedure.
2.7. EBSD sample preparation
The second half of each compressed sample was electrolytically polished for the EBSD analysis using standard metallographic techniques, finishing with vibrational silica polishing to a finish of 0.05 µm. The samples were then jet polished at a temperature
between 8 ˚C and 18 °C in a 3 mol L-1 solution of sulfuric acid in
ethanol. The electrolytic polishing voltages, currents and times were 30-40 V, 1-2A, and approximately 30 s, respectively.
2.8. Electron Backscatter Diffraction Analysis (EBSD)
EBSD was conducted using a Zeiss Sigma field emission gun scanning electron microscope equipped with a “Nordlys F” EBSD detector [27]. The EBSD data collection and subsequent indexing was performed using the AZtec 3.3 and 4.3 softwares. In all five Supplements I-V, the EBSD data acquisition was
performed by indexing the face-centred cubic (γ) phase. For each
condition, at least two EBSD scans were acquired. From the data, several microstructural parameters were evaluated using the Channel5 software “TANGO” [27] (Supplement I) and the software TSL OIM Analysis 7 (supplements II, III, IV and V).
53
In the present work, a grain boundary is defined by EBSD when
the change in orientation across it exceeds 5° [21,28]. This is
known as the grain tolerance angle [29]. Twin boundaries
(defined as boundaries with a misorientation of 60° about 〈111〉
axes) were ignored when estimating the recrystallized grain size. Boundaries which are classified as high-angle (misorientation
angle, * > 10°) are interpreted as fully formed grain boundaries,
whereas low angle boundaries (2° < * < 10°) are interpreted
as sub-grain boundaries comprising a high density of
dislocations. The minimum misorientation angle, *, detected
between grains was 2° and the maximum was 62.7°. The grain
orientation spread technique was used to distinguish the dynamically recrystallized grains from the deformed matrix [30,31]. In this technique, the average difference in orientation between each data acquisition point and the mean for the entire
grain is calculated. If a grain has a grain orientation spread of <
2°, it is interpreted as a recrystallized grain. Higher values are taken to indicate that a grain has either not undergone recrystallization or has deformed significantly after being formed by recrystallization [30,32–38]. This is more reliable than using the grain size to identify recrystallized grains [34]. The threshold value for classifying grains as recrystallized or deformed is given variously in literature as 1-2°, 3° or 5° [33–
37]. During this study, a value of 2° was found to be appropriate.
The coincident site lattice model was also employed to describe
the grain boundary characteristics, in which one in every Σ
atoms are coincident in the two grains across the boundary. The
Brandon criterion, i.e. tolerance, *m = _…„†b = 15°Σ ⁄c [38],
where K = 15° and & = 0.5 , was used to assess whether a
boundary is a coincident site lattice boundary. Sigma (Σ) can
have values between 1 to 49. High angle grain boundaries
54
random (general) boundaries, Σ > 29. Thereafter, the TSL OIM
Analyzer software was also used to identify twin boundaries in order to exclude these from the grain size calculations (for twin
boundaries, the misorientation angle, * = 60° with a tolerance
of 5° around the theoretical <111> axes).
EBSD was also used to investigate deformation texture and recrystallized texture by the inspection of the inverse pole figure of each sample. The fraction of recrystallized grains, the number density and size of dynamically recrystallized grains (excluding twins) were investigated. Data were cleaned before analysis using grain dilation to minimize the influence of noise on the results. A minimum grain size of 10 pixels was also defined. Table 7 shows a step size, a covering area and a threshold of recrystallized grain for each supplement.
Table 7. A scanning of step size, covering area and grain orientation spread
Experiment Step size,
[µm] Area, [mm2] Recrystallized grain Supplement I 1.5, 0.5 ~1.0, ~ 0.11 GOS <2° Supplement II 3.0, 1.0 ~4.03, ~0.253 GOS <2°
Supplement III 2.0 ~4.05 GOS ≤ 1.3°
Supplement IV 3.0, 0.5,
0.75 ~4.05, ~ 0.11,
~ 4.03
GOS ≤ 1.0°
