Materials Processing and Texture Ceramic Transactions, Volume 200 A Collection of Papers Presented at the 15th International Conference on Textures of Materials (ICOTOM 15) June 1-6, 2008 Pittsburgh, Pennsylvania
Edited by A. D. Rollett
The American Ceramic Society A À
WILEY
A John Wiley & Sons, Inc., Publication
This Page Intentionally Left Blank
Materials Processing and Texture
This Page Intentionally Left Blank
Materials Processing and Texture Ceramic Transactions, Volume 200 A Collection of Papers Presented at the 15th International Conference on Textures of Materials (ICOTOM 15) June 1-6, 2008 Pittsburgh, Pennsylvania
Edited by A. D. Rollett
The American Ceramic Society A À
WILEY
A John Wiley & Sons, Inc., Publication
Copyright © 2009 by The American Ceramic Society. All rights reserved. Published by John Wiley & Sons, Inc., Hoboken, New Jersey. Published simultaneously in Canada. No part of this publication may be reproduced, stored in a retrieval system, or transmitted in any form or by any means, electronic, mechanical, photocopying, recording, scanning, or otherwise, except as permitted under Section 107 or 108 of the 1976 United States Copyright Act, without either the prior written permission of the Publisher, or authorization through payment of the appropriate per-copy fee to the Copyright Clearance Center, Inc., 222 Rosewood Drive, Danvers, MA 01923, (978) 750-8400, fax (978) 750-4470, or on the web at www.copyright.com. Requests to the Publisher for permission should be addressed to the Permissions Department, John Wiley & Sons, Inc., 111 River Street, Hoboken, NJ 07030. (201) 748-6011, fax (201 ) 748-6008, or online at http://www.wiley.com/go/permission. Limit of Liability/Disclaimer of Warranty: While the publisher and author have used their best efforts in preparing this book, they make no representations or warranties with respect to the accuracy or completeness of the contents of this book and specifically disclaim any implied warranties of merchantability or fitness for a particular purpose. No warranty may be created or extended by sales representatives or written sales materials. The advice and strategies contained herein may not be suitable for your situation. You should consult with a professional where appropriate. Neither the publisher nor author shall be liable for any loss of profit or any other commercial damages, including but not limited to special, incidental, consequential, or other damages. For general information on our other products and services or for technical support, please contact our Customer Care Department within the United States at (800) 762-2974, outside the United States at (317) 572-3993 or fax (317) 572-4002. Wiley also publishes its books in a variety of electronic formats. Some content that appears in print may not be available in electronic format. For information about Wiley products, visit our web site at www.wiley.com.
Library of Congress Cataloging-in-Publication ISBN 978-0-470-40834-6 Printed in the United States of America. 10
9 8 7 6 5 4 3 2 1
Data is available.
Contents
Preface Acknowledgments
xv XVII
FRICTION STIR WELDING AND PROCESSING Texture Evolution during Friction Stir Welding of Stainless Steel
3
Jaehyung Cho
Texture Development in Aluminum Friction Stir Welds
17
R i c h a r d W . F o n d a , Keith E. Knipling, J o h n F. Bingert, A n t h o n y P. R e y n o l d s , W e i T a n g , Kevin J . C o l l i g a n , a n d J o h n A. W e r t
Microstructural Studies of a Friction-Stir-Welded AA5052
29
N . T . K u m b h a r , S a n t o s h . K. S a h o o , R. T e w a r i , G . K. D e y , I. S a m a j d a r , a n d K. B h a n u m u r t h y
Effect of Microstructure on Mechanical Properties of Friction-Welded Joints between Aluminum Alloys (6061, 5052) and 304 Stainless Steel
35
Ν . T . K u m b h a r , A. Laik, G . K. D e y , Κ. B h a n u m u r t h y , J . Krishnan, D. J . D e r o s e , R. L. S u t h a r , S . K. S a h o o , a n d I. S a m a j d a r
Effect of Texture on Fracture Limit Strain in Friction Stir Processed AZ31F3Mg Alloy
43
Y u t a k a S . S a t o , Akihiro S u g i m o t o , Hiroyuki K o k a w a , a n d C h a n g - W o o L e e
Microstructure and Texture Analysis of Friction Stir Welds of Copper
53
T . S a u k k o n e n . , K. S a v o l a i n e n , J . M o n o n e n , a n d H. H ä n n i n e n
TEXTURE AND ANISOTROPY IN STEELS The Effects of Texture on the Corrosion Resistance of SS304 Stainless Steel
63
M . Azzi a n d J . A. S z p u n a r
ν
Microstructural Characterization of Dual Phase Steels by Means of Electron Microscopy
71
L i e s b e t h B a r b é a n d Kim V e r b e k e n
Grain Size Effect on the Tensile Behavior of a FeMnC TWIP Steel in Relation with the Microstructure and Texture Evolutions
79
D a v i d Barbier, N a t h a l i e G e y , Alain S é b a s t i e n , a n d M i c h e l H u m b e r t
Combination of Two X-Ray Diffraction Settings for Determining the Texture of Small Sheared Samples
87
D. Barbier, Β. Bolle, J . J . F u n d e n b e r g e r , C . Laruelle, a n d A. T i d u
Texture of Non-Oriented Rolled Steels by X-Ray and Electron Diffraction
95
M a r t i n Öernik, M i l o s P r e d m e r s k y , A n d r e j L e s k o , a n d Erik Hilinski
Effect of Coating Texture on Powdering Behavior of Industrially Produced Galvannealed Coating on IF, IFHS and HSQ Grade Steels
103
A. C h a k r a b o r t y , R.K. R a y , a n d D. B h a t t a c h a r j e e
Including Elastical Anisotropy Due to Texture in a Finite Element Model for the Prediction of Residual Stresses in Welds
115
Koen Decroos, Carsten Ohms, and Yvan Houbaert
Texture and Microstructure Evolution in a Fe-Si CGO Sheet during the Processing Route before Secondary Recrystallization
123
F r a n c i s c o C r u z - G a n d a r i l l a , R i c h a r d P e n e l l e , Thierry B a u d i n , H e c t o r M e n d o z a Leon, and J. G. C a b a n a s - M o r e n o
TEM Study of the Inhibition Layer of Commercial Hot-Dip Galvanized Steels
131
J a c k Elia, Ε. Petit, J - S L e c o m t e , B. G a y , a n d V . P i t c h o n
Measurement of the Texture Sharpness in Grain-Oriented Electrical Steels
143
M . F r o m m e r t , C . Z o b r i s t , D. R a a b e , S . Z a e f f e r e r , a n d L. L a h n , Α. B ö t t c h e r
Effects of Composition and Coiling Temperature on Precipitation and Texture Formation in a Few Interstitial Free High Strength Steels
151
P a m p a Ghosh and R Κ Ray
Microstructure and Texture Development in Warm Rolled High Strength IF Steels Containing Cu, Ρ and Β
161
Arunansu Haldar and Chiradeep Ghosh
Effect of Tin on the Texture of Nonoriented Electrical Steels C h u n - K a n H o u , T o n g - H a n L e e , a n d C h u n - C h i h Liao
vi
·
Materials Processing and Texture
173
Effect of Cold Rolled Reduction on the Texture and Plastic Strain Ratio of 2205 Duplex Stainless Steel
183
C h u n - K a n H o u a n d C h a n g - C h i a Liu
Importance of Dislocation Substructures in Texture Control of Low Carbon Steel
193
H i r o s u k e Inagaki
Texture Control in Manufacturing Current and Future Grades of Low-Carbon Steel Sheet
207
Leo A.I. Kestens and R o u m e n Petrov
Relative Preference for Strain Localizations in Ultra Low Carbon Steel
217
R. Khatirkar, L.A.I. K e s t e n s , R. Petrov, a n d I. S a m a j d a r
Influence of Texture on Surface-Roughening during Press Forming in Ti Added High Purity Ferritic Stainless Steel Sheets
223
K e n K i m u r a , M a s a h a r u H a t a n o , a n d Akihiko T a k a h a s h i
Effect of Tin on the Recrystallization Behavior and Texture of Grain Oriented Electrical Steels
231
C h u n - C h i h Liao, C h u n - K a n H o u
EBSD Investigation of Phase Transformation in Low Carbon Steels
243
S o p h i e Lubin, A n n e - F r a n ç o i s e G o u r g u e s - L o r e n z o n , H é l è n e R e g l e , Frank M o n t h e i l l e t , a n d Brigitte B a c r o i x
A Monte-Carlo Algorithm for the Zener Drag Effect in the Fe-3%Si Magnetic Alloys
251
N . M a a z i , R. P e n e l l e , A. L. Etter, a n d T. B a u d i n
Effect of Strain and Strain Path on Deformation Twinning and Strain Induced Martensite in AISI 316L and 304L Austenitic Stainless Steel
257
S . K. M i s h r a , P. P a n t , K. N a r a s i m h a n a n d I. S a m a j d a r
Analysis of the Formation and Growth of Recrystallizing Grains in IF Steels by In-Situ EBSP
265
H. N a k a m i c h i a n d F. J . H u m p h r e y s
Stability of Secondary Recrystallization in Grain-Oriented Silicon Steel
275
Eiichi N a m b a , S a t o s h i Arai, H o t a k a H o m m a , a n d Y o s h i y u k i U s h i g a m i
A Practical Investigation into Identifying and Differentiating Phases in Steel using Electron Backscatter Diffraction
285
M a t t h e w M . N o w e l l , Stuart I. W r i g h t , a n d J o h n O . C a r p e n t e r
Materials Processing and Texture
vii
Influence of Precipitate Distribution on Normal Grain Growth Simulation at the Beginning of the Secondary Recrystallization Annealing in a Fe-3%Si Electrical Steel HiB - Role of the MnS and AIN Precipitates
293
C . S e k k a k , N . R o u a g , a n d R. P e n e l l e
Microstructure and Texture Development during Cold Ring Rolling of 100Cr6
301
K. R y t t b e r g , M . K n u t s o n W e d e l , L. N y b o r g , a n d P. D a h l m a n
Development of Texture, Microstructure and Grain Boundary Character Distribution after Heavy Cold Rolling and Annealing in a Boron Added Interstitial Free (IF) Steel
311
Rajib S a n a a n d R. K. R a y
Texture of Nb-Containing Ferritic Stainless Steel after Secondary Recrystallization
317
R o d r i g o P. S i q u e i r a , H u g o R. Z . S a n d i m , a n d T a r c i s i o R. Oliveira
Quantitative Prediction of Transformation Texture in Hot-Rolled Steel Sheets by Multiple KS Relation
325
T o s h i r o T o m i d a , M a s a y u k i W a k i t a , M i t s u r u Y o s h i d a , a n d Norio Imai
Crystallographic Characterization of a Phosphorus Added TRIP Steel
333
Kim V e r b e k e n , L i e s b e t h B a r b é , a n d Dierk R a a b e
Microstructural and Textural Evolutions in a Cold Rolled High Mn TWIP Steel
341
K i m V e r b e k e n , Lieven B r a c k e , L e o K e s t e n s , a n d J a n P e n n i n g
Stored Energy Evolution in a Cold-Rolled IF Steel
349
A. W a u t h i e r , B. B a c r o i x , T . C h a u v e a u , O . C a s t e l n a u , a n d H. R é g l é
A Comparison of Texture Measurements Via EBSD and X-RAY
357
S t u a r t I. W r i g h t a n d M a t t h e w M . N o w e l l
The Measurement Approach of X-Ray Residual Stress in Fe-Ni Alloy with Different Textures
365
Qiulin W u , X i a o j u n H u , Z u m i n g F u , H a i y a n Z h a o , Bing Z h a n g , Y o n g W a n g , W e i F a n g , a n d Limin Li
Influence of Grain Orientation on Magnetic Aging Behavior of a NonOriented Silicon Steel Sheet L. X u , W . M a o , P. Y a n g , a n d H. F e n g
viii
·
Materials Processing and Texture
371
EFFECTS OF MAGNETIC FIELDS Eutectoid Point Shift and Orientation Relationships between Ferrite and Cementite in Pearlite in a High Magnetic Field
381
Y . D. Z h a n g , C . Esling, H. Klein, X . Z h a o , a n d L. Z u o
Characteristics of Recrystallization Texture of Cold-Rolled IF Steel Sheet Annealed with a Magnetic Field in the Transverse Direction
389
Y a n W u , C h a n g s h u H e , Y u d o n g Z h a n g , X i a n g Z h a o , Liang Z u o , a n d C l a u d e Esling
Crystallographic Features during Martensitic Transformation in Ni-Mn-Ga Ferromagnetic Shape Memory Alloys
397
D. Y . C o n g , Y . D. Z h a n g , C . Esling, Y . D. W a n g , X . Z h a o , a n d L. Z u o
Grain Boundary Misorientation and CSL in Magnetically Annealed Fe-0.75%Si
405
C . M . B . Bacaltchuk, G.A. Castello-Branco, and H. Garmestani
The Effect of Ultra-High Magnetic Fields on Grain Growth and Recrystallization in Metals
413
A . D . S h e i k h - A l i a n d H. G a r m e s t a n i
New Approach to Texturing and Grain Boundary Engineering by Magnetic Field Application
421
T a d a o W a t a n a b e , S a d a h i r o T s u r e k a w a , Y u d o n g Z h a n g , X i a n g Z h a o , Liang Z u o , a n d C l a u d e Esling
Effects of Magnetic Field Direction on Texture Evolution in a Cold-Rolled IF Steel Sheet during High Magnetic Field Annealing
435
Y a n W u , C h a n g s h u H e , X i a n g Z h a o , a n d Liang Z u o
HEXAGONAL METALS Electrochemical Behavior of (001), (100) and (110) Ti Single Crystals under Simulated Body Fluid Condition
443
M . A z z i , S . Faghihi, M . T a b r i z i a n , a n d J . A. S z p u n a r
Identification and Rationalization of Secondary Twin Variants in a Magnesium Alloy
451
M . R. B a r n e t t a n d P. C i z e k
Textures in HCP Titanium and Zirconium: Influence of Twinning
461
Nathalie Bozzolo and Francis W a g n e r
Texture and Microstructure Evolution during Asymmetric Rolling of Mg Alloys
473
Jaehyung Cho, Hyung-Wuk Kim, and Suk-Bong Kang
Materials Processing and Texture
•
ix
Prediction of Anisotropic Properties in Mg Alloy Sheets Using the Crystal Plasticity Finite Element Method
483
S . - H . C h o i , J.K. C h o i , H . W . L e e , a n d D . H . Kim
Strain Hardening Behaviour of an Initially Textured TI6AL4V Titanium Alloy as a Function of Strain Rate and Compression Direction
491
Frederik C o g h e , Luc R a b e t , a n d P a u l V a n H o u t t e
Texture Evolution during the Static Recrystallization of Three Binary Mg-Y Alloys
501
R. C o t t a m , J . R o b s o n , G . Lorimer, a n d B. D a v i s
Grain Size and Orientation Distribution Function of High Purity α-Titanium
509
B r a d l e y S . F r o m m , B r e n t L. A d a m s , S a d e g h A h m a d i , a n d M a r k o K n e z e v i c
Textures in Titanium Alloys - An Industrial Perspective on Deformation, Transformation and Properties
521
D a v i d R u g g , D a v i d Furrer, a n d Nigel B r e w i t t
Annealing Related Microstructural Developments in a Two-Phase ZR-2.5 Nb Alloy
533
V . D . H i w a r k a r , S . K . S a h o o , I. S a m a j d a r , K. N a r a s i m h a n , Κ. V . M a n i krishna, G.K. D e y , D. S r i v a s t a v , R. T i w a r i , a n d S . B a n a r j e e
Texture Aspects of Delayed Hydride Cracking in Products from Zr-Based Alloys
539
M a r g a r i t a I s a e n k o v a a n d Yuriy Perlovich
Effect of Twinning on the Strain Hardening Behaviors of Two Mg Alloys Deformed Along Four Different Strain Paths
547
L. J i a n g a n d J . J . J o n a s
EBSD Study of Annealing Rolled Zirconium
555
L. J i a n g , M . T . P é r e z - P r a d , O.A. R u a n , a n d M . E . K a s s n e r
Texture Evolution during Multi-Pass Equal Channel Angular Extrusion of Beryllium
563
Saiyi Li, D a v i d J . A l e x a n d e r , Irene J . B e y e r l e i n , a n d D o n a l d W . B r o w n
A Mechanism of Determining the Formability of AZ31 near Room Temperature
571
H u a l o n g Li, Emilie H s u , a n d J e r z y S z p u n a r
An EBSD Study of the Misorientations Related to Dynamic Recrystallization in Mg AM30 Deformed At High Temperatures Etienne Martin, S t é p h a n e G o d e t , Lan Jiang, A b d e l b a s e t Elwazri, Pascal J. Jacques, and John J. Jonas
χ
·
Materials Processing a n d Texture
577
Texture Evolution in Boron Modified Ti-6AI-4V Alloy
585
S h i b a y a n R o y , Nilesh G u r a o , S a t y a m S u w a s , S . T a m i r i s a k a n d a l a , R. S r i n i v a s a n , a n d D.B. Miracle
Heterogeneous Deformation in Single-Phase Zircaloy 2
593
S . Κ. S a h o o , V. D. H i w a r k a r , I. S a m a j d a r , P. P a n t , D. Srivastav, R. T e w a r i , G . K. D e y a n d S . B a n e r j e e
Texture of Magnesium Alloy Sheets Heavily Rolled by High Speed Warm Rolling
601
T e t s u o S a k a i , S a t o s h i M i n a m i g u c h i , Hiroaki K o h , a n d Hiroshi U t s u n o m i y a
Microstructure and Texture Gradient in Titanium Severely Strained by Friction Roll Processing and Subsequent Annealing
609
Meiqin Shi, Tomohiro U m e t s u , Yoshimasa T a k a y a m a , Hajime Kato, and Hideo Watanabe
Evolution of Transformation Texture in a Metastable ß-Titanium Alloy
617
S a t y a m S u w a s , Nilesh P. G u r a o , A s h k a r Ali
Dislocations in Textured Mg-4.5%AI-1 %Zn Alloy: TEM and Statistical Studies
629
V . N . T i m o f e e v , V . N . S e r e b r y a n y , a n d Y u A. Z a l i z n y a k
Microstructure and Micro-Texture Evolution of Compression Twins in Magnesium
637
P. Y a n g , L. M e n g , Χ . Li, W . M a o , a n d L. C h e n
TEXTURE IN MATERIALS DESIGN Accessing the Texture Hull and Properties Closure by Rotation and Lamination: Results in the Primitive Basis of Dilation Functions
647
Brent L. A d a m s a n d D a v i d T. F u l l w o o d
Correlation Relationship between Yield Strength and the Elastic Compliance for the Polycrystalline Nickel 201
657
S a d e g h A h m a d i , Brent L. A d a m s , D a v i d T . F u l l w o o d , a n d B r a d l e y S. F r o m m
Modeling Rolling Texture of Twinned Cu Single Crystals
665
Khaled Al-Fadhalah, Armand Beaudoin, and Marek Niewczas
Multiscale Modeling of Equal Channel Angular Extruded Aluminum with Strain Gradient Crystal Plasticity and Phenomenological Models
671
L. D u c h ê n e , M . G . D . G e e r s , W . A . M . B r e k e l m a n s , E. C h e n , B. V e r l i n d e n , a n d A.M. Habraken
Materials Processing and Texture
·
xi
Effect of Deformation Constraints on the Texture Formation in Al5mass%Mg Solid Solution at High Temperatures
679
K a z u t o O k a y a s u , M a s a y u k i S a k a k i b a r a , a n d Hiroshi F u k u t o m i
Spectral Methods in the Statistical Description and Design of Microstructure
687
D.T. F u l l w o o d , B.L. A d a m s , K.A. S t e v e n s , S.R. N i e z g o d a , a n d S . R . Kalidindi
A Comparison of Deformation Textures and Mechanical Properties Predicted by Different Crystal Plasticity Codes
701
C r a i g S . H a r t l e y , Paul R. D a w s o n , D o n a l d E. B o y c e , S u r y a R. Kalidindi, M a r k o K n e z e v i c , C a r l o s T o m é , R i c a r d o L e b e n s o h n , S . L e e S e m i a t i n , T o d d J . Turner, a n d A y m a n A. S a l e m
Simulation of Texture Development in Pure Aluminum Deformed by Equal Channel Angular Pressing
713
M a j i d H o s e i n i , M a h m o o d M e r a t i a n , H u a l o n g Li, a n d J e r z y S z p u n a r
Plastic Heterogeneity Due to Grain Boundaries and Its Influence on Global Deformation Textures
721
A n a n d Κ Kanjarla, Paul V a n H o u t t e , a n d Laurent D e l a n n a y
On the Correlation of Surface Texture and Strain Induced Surface Roughness in AA6XXX Aluminum Sheet
731
S . K ü s t e r s , M . S e e f e l d t , a n d P. V a n H o u t t e
The Variation Range for the Results of FCC Texture Simulations Based on {111 }<110> Slip
743
T o r b e n Letters
Development of Low Magnetism and Strongly Cube-Textured Composite Ni-7%W/Ni-10%W Substrate for Coated Superconductor
751
D a n - m i n Liu, Fei H a o , J i u - x i n g Z h a n g , Y a n - c a o H u , M e i - l i n g Z h o u , a n d W e i - p e n g Liu
Two-Point Orientation Coherence Functions: An Orthonormal Series Solution
759
P e t e r R. M o r r i s
The Influence of Shear Bands on Microtexture Evolution in Polycrystalline Copper
765
H. Paul a n d J . H . Driver
Influence of Dislocation Stress on the Corrosion Behavior of (111) and (123) Surface Plane of Aluminum Single Crystal Dashi Pei, W e i m i n M a o , a n d Huiping Feng
xii
·
Materials Processing and Texture
775
Texture Randomization by Intense Shearing in Layered AI and AI (0.3%Sc) ARB-Processed Sheet
783
M . Z . Q u a d i r , O . A l - B u h a m a d , L. B a s s m a n , a n d M . Ferry
Texture Development of Molybdenum Sheets during Last Step of Heat Treatment
791
C . - G . O e r r e l , I. H ü n s c h e , W . S k r o t z k i , A. Lorich, a n d W . K n a b l
Solute, Superplasticity and Texture Change in Aluminium Alloys
801
K. S o t o u d e h , P . S . B a t e , a n d F.J. H u m p h r e y s
Texture-Based Plastic Potentials in Stress Space
809
Albert V a n B a e l , S a m p a t h K. Y e r r a , Paul V a n H o u t t e , a n d Albert V a n Bael
Hierarchical Multi-Level Modelling of Plastic Anisotropy using Convex Plastic Potentials
817
Paul V a n H o u t t e , S a m p a t h K. Y e r r a , a n d Albert V a n Bael
Author Index
827
Materials Processing and Texture
xiii
This Page Intentionally Left Blank
Preface
It has been a pleasure, o f a rather hard working type, to edit the proceedings for the 15th International Conference on Textures o f Materials (ICOTOM 15). First and foremost I acknowledge the hard work and time that the symposium organizers have devoted to this task. This represents truly a community effort to sustain the development o f this field o f materials research. Texture, or crystallographic orientation, plays an important role in many different areas o f materials research and development. It grew out o f metallurgical processing but now is significant to geology, welding, thin films, electronic devices, fuel cells, biomaterials and other applications too numerous to list. I am pleased that this meeting brings together such a broad range o f interests. The papers included in this Ceramic Transactions volume (as well as its companion volume—Ceramic Transactions V o l u m e 2 0 1 , Applications o f Texture Analysis) illustrate this broad range o f applicability o f the tools. Friction Stir Welding has undergone extensive development in recent years so it is appropriate that this volume has an entire section on the topic. Texture development and its effect on properties in steels remains of great interest to the community. Hexagonal metals such as titanium have greatly increased their range o f applications so it is not surprising that texture is an essential aspect o f these complex alloy systems. Magnetic fields have a strong effect on many materials and their influence on processing and resulting textures has attracted substantial interest. Lastly, materials design is being placed on a much firmer theoretical and computational foundation, as is evident in the work reported here. I have no doubt that the community will find these proceedings as essential in their work as they have for all the previous ICOTOM conferences. I would like to thank the many colleagues w h o helped in organizing the conference and especially Dr. Carlos Tomé o f Los Alamos who put together the workshops that are such a distinctive feature o f this conference. Many o f these individuals also made substantial contributions to the review process o f the proceedings. Their names are listed on the Acknowledgments page. I also thank the National Science Foundation, the Air Force Office o f Scientific Research and the Alcoa Technical Center for their support o f the conference. Their
XV
sponsorship made it possible for the many young scientists and engineers to attend the conference. The staff o f The American Ceramic Society (ACerS) were unfailingly supportive and professional in their management o f the conference. The Minerals, Metals and Materials Society (TMS) also played an important role in developing the workshops. I thank the many students and staff w h o contributed to the smoothing running and organization o f the meeting. A.D.
TONY
Carnegie
xvi
·
ROLLETT
Mellon
University
Materials Processing and Texture
Acknowledgments
Conference proceedings co-editors Brent Adams, Brigham Young University, U S A Katayun Barmak, Carnegie Mellon University, U S A Thomas Bieler, Michigan State University, U S A Keith Bowman, Purdue University, U S A Paul Dawson, Cornell University, U S A Roger Doherty, Drexel University, U S A OlafEngler, Hydro Aluminium Deutschland GmbH and R & D Center, Germany Hamid Garmestani, Georgia Institute of Technology, U S A Amit Goyal, Oak Ridge National Laboratory, U S A Erik Hilinski, U S Steel, U S A Roland Loge, The Centre for Materials Forming, France S. Lee Semiatin, Air Force Research Laboratory, U S A Robert Suter, Carnegie Mellon University, U S A Carlos Tomé, Los Alamos National Laboratory, U S A Sven Vogel, Los A l a m o s National Laboratory, U S A Hasso Weiland, Alcoa Technical Center, U S A Stuart Wright, E D A X - T S L , U S A
xvii
This Page Intentionally Left Blank
Friction Stir Welding and Processing
This Page Intentionally Left Blank
T E X T U R E E V O L U T I O N D Ü R I N G FRICTION STIR W E L D I N G O F STAINLESS STEEL Jaehyung Cho Korea Institute of Materials Science, 66 Sangnam-dong, Changwon-city, Kyungnam, South Korea 641-010 Email: j h c h o @ kims.re.kr ABSTRACT Texture evolution during ruction stir welding (FSW) of stainless steel was investigated using both experimental measurements and model predictions based on a polycrystal plasticity. Material undergoes heating and deformation sufficient to alter its microstructural state during FSW. Texture evolution is quite complex in FSW, owing in particular to its dynamics under the induced stirring motion. Modeling approaches help to better understand h o w the stirring can both strengthen and weaken textures, which is an essential element in a complete description of texturing during FSW. In a two-dimensional way, texture stability is addressed from the computed velocity gradients along streamlines of the flow field. The texture is assumed to be uniform initially and shows monoclinic sample symmetry after deformation. Upstream and downstream of the tool, the deformation is nearly monotonie, causing little change of the texture. Around the tool pin, the texture strengthens, weakens, and restrengthens. The mixture of pure and simple shear textures was found. The effects of frictional conditions with the tool pin and shoulder on the complicated flow and texture evolution in the through thickness were examined with a three-dimensional approach. Trends in regard to strengthening and weakening of the texture were discussed in terms of the relative magnitudes of the deformation rate and spin. The computed textures were compared to E B S D measurements and were discussed with respect to distributions along orientation fibers and the dominant texture component along the fibers. 1.
INTRODUCTION
Material undergoes severe heating and deformation during friction stir welding (FSW) and its microstructural features, such as texture, grain size and shape changed. Material in the weld zone experiences strains large enough to drive the texture toward a few dominant components. Dynamic recrystallization is found, particularly in aluminum alloys because of the elevated temperatures present around the tool. Materials flow is so complicated and texture evolution is also dynamic and complex. A number of studies characterizing the texture alterations and variations of the mechanical responses have been reported [1]. From electron backscatter diffraction (EBSD) measurements of various aluminum alloys, it has been shown that severe texture gradients exist through the thickness direction and across the width of parts joined by FSW. Shear-type texture components, which are rotated from the ideal shear components [2], often are observed. From a microtexture analysis after postweld heat treatment, dynamic recrystallization during F S W generates the recrystallized grains at the weld center instead of static recrystallization [3]. Using a stop-action method, the initial stages of texture evolution ahead of the pin in Al-Li 2195 alloys have been examined [4, 5]. The primary mechanism of grain refinement appeared to be grain subdivision induced by deformation, although the initial grain refinement seemed to be followed by continuous dynamic recrystallization and recovery [4]. The fine subgrains ahead of the tool display a typical fee shear texture after suitable rotation. In addition to grain subdivision, geometric dynamic recrystallization was suggested to contribute the ultrafined grain structure with shear orientations in the nugget [5]. The recrystallization phenomenon during F S W of 304L stainless steel was examined in Ref. [6]. The dynamically recrystallized grains have two types of orientation distributions, i.e., textures with one dominant component and distributions along the adjacent fiber. Texture evolution can be computed by coupling the finite element solution for the material
3
T e x t u r e Evolution d u r i n g Friction Stir W e l d i n g of S t a i n l e s s S t e e l
motion with crystal plasticity for texture analysis. Crystal flow over orientation space is driven by a reorientation velocity field that smoothly varies. Some orientations are stable under a monotonie deformation. While, there are no stable components for torsion textures even though the deformation remains monotonie to large strains. The relationship between the derivatives of the reorientation velocity at ideal orientations and the local flows around the orientations was explored in [7], The dynamical behavior of crystal flows was analyzed for idealized planar crystals in [8]. Consideration of flow dynamics also has been reported in the case of a flow field that does not exhibit equilibria. The continuous flow associated with simple shear (torsion) in terms of the structure of the reorientation velocity field was investigated in [9, 10]. An analysis of the stability of texture components in fee systems under simple and pure shear modes of deformation was also presented in [11]. The focus of this paper is on modeling texture evolution in 304L stainless steel as it is altered by the F S W process and on comparing the computed textures to ones measured using E B S D . The material response during FSW was modeled using an isotropic, strain hardening, viscoplastic constitutive equation in a strongly coupled thermomechanica! formulation. The model equations were solved in an Eulerian reference frame for both two-dimensional and two-dimensional geometries using a finite element method [12, 13]. Temperature and strength distributions from the simulations were compared with experimental results. Based on the material response in two-dimensional model, studies on the cry stallographic texture evolution focused on an examination of the dynamics of the texture evolution. It was found that some combinations of the symmetric and skew parts of the velocity gradient induced constantly changing texture patterns, while others promoted monotonie texturing. The effect of friction between the tool and the workpiece on the three dimensionality of the flow was investigated and its consequences regarding texture evolution was discussed. It is considered how the material rotation impacts the strength of the texture and the occurrence of fiber textures and strong texture components. Finally, predicted textures were compared with ones measured by E B S D over a cross section of the welded plate. A three-dimensional schematic diagram illustrating the F S W process is shown in Fig. 1.
Figure 1. A schematic diagram of the friction stir welding process. 2.
MODEL GEOMETRY AND BOUNDARY CONDITIONS The complete F S W process includes the plunge, dwelling, main welding, and pull-out operations. For simplicity, a stress was placed on the main welding process during successful steady-state welding. During the welding, the effects of the tool pin and shoulder produce complex material motions. More details on two- and three-dimensional modeling are found in references [12, 13].
4
Materials Processing and Texture
T e x t u r e E v o l u t i o n d u r i n g Friction Stir W e l d i n g of S t a i n l e s s S t e e l
Figs. 2(a) and 2(b) display two- and three-dimensional discretized meshes of FSW. The three-dimensional model (Fig. 2b) has both the tool shoulder (surface 6) and tool pin (surface 5), while in the two-dimensional model (Fig. 2a) only the tool pin is considered. Elements are concentrated near the tool surface where gradients of the velocity and temperature are greatest. The tool pin is located in the center of the region and material moves from left to right. The tool spins clockwise with an angular velocity, Ω, of 5rad/s. Therefore, the lower part of the region corresponds to the advancing side and the upper part to the retreating side. Boundary conditions and model geometries are listed in Table 1. Friction boundary conditions were used to represent the contact condition between the workpiece and the tool pin and shoulder. On the shoulder, the friction coefficients were allowed to linearly decrease with radial distance. On the tool pin, the friction coefficients were constant, but an axial component of traction was added to simulate the pin thread effect. The ratio of the axial component to the tangential component determines the thread angle. It is known that the tool shoulder generates m o r e heat than the tool pin. Even with the lower friction coefficient in these simulations, the tool shoulder effect remains dominant in deformation and heating. Table I. Boundary conditions for F S W models Translation Rotation Pin angle Shoulder Type iT, [mmIs] Ω [radis] [Deg] On/off 2D 3D
1.7 1.0
5 5
N/A 5
Geometry Pm radius[mm]
Geometry Shoulder radius [mm]
to 3
N/A 12
N/A on
Note: u is the welding speed; Ω is the angular velocity of the tool pin. Friction coefficients for the shoulder s
(ßühmiidcr)
3
linearly decrease with distance from the tool pin, from 150x10* to 150x10° [N-s/m ]. Friction 9
3
coefficients for the pin(ßmn) have constant values of 150x10 [N-s/m ].
(a) (b) Figure 2. Finite element meshes and overall descriptions of the surface boundary conditions for FSW models. In Fig. 2(a) (2D): Inlet; . O u t l e t ; _ Top and Bottom (adiabatic); _ Tool pin. In Fig. 2(b) (3D): Inlet; and T o p and Bottom (adiabatic); Outlet; Tool pin; Tool shoulder; Plate to be welded (adiabatic).
Materials Processing and Texture
·
5
T e x t u r e Evolution d u r i n g Friction Stir W e l d i n g of S t a i n l e s s S t e e l
3. M O D E L I N G T E X T U R E E V O L U T I O N The F S W tool usually consists of two parts, the shoulder and pin. Both pin and shoulder contribute to complex, three-dimensional material flow patterns. Visualization techniques have been employed to study the complicated material flow experimentally, as reported in [14, 15]. Simulation of the process also offers the opportunity to better understand the effects of these aspects of the tool geometry on resulting flow and the resulting evolution of microstructure. The simulations presented here utilize a methodology that decouples the computations for the velocity and temperature fields from the integration of the texture evolution equations. The procedure consists of three basic steps. The first step is to compute the velocity and temperature fields over the control volume. This is followed by a step to extract thermomechanical histories along streamlines leading to target positions at the outlet. The final step is to integrate evolution equations for texture along the streamlines using the thermomechanical data. In using this decoupled procedure, it is assumed that the plastic anisotropy present in the yield surface due to crystallographic texture has little influence on the overall flow field. Isotropic strain hardening or softening was allowd, so the overall magnitude of the flow stress does change as material is deformed and heated. Full coupling is possible for other applications, but would require considerably greater execution times for the same spatial resolution [ 1 6 , 1 7 ] . In this study, greater spatial resolution over full coupling was preferred. 3.1 C O M P U T I N G T H E E V O L U T I O N O F C R Y S T A L L O G R A P H I C T E X T U R E Texture evolution is computed using a framework based on polycrystal plasticity [18] together with a methodology of integrating the texture along the streamlines. The macroscopic velocity of a representative volume of material at a specified spatial position χ in an Eulerian flow field is n. The velocity gradient L is obtained by taking the spatial gradient of the velocity field L = 3 u / 3 x . This motion at the macroscopic scale is linked to the crystal scale by adopting the extended Taylor hypothesis such that all the crystals comprising the representative volume experience the macroscopic velocity gradient [19]. In the case of the high symmetry crystals and large strains considered here, this is a reasonable approximation. As our attention is focused on regions that experience large deformations, elastic effects are neglected and crystals are assumed to deform solely by crystallographic slip. In this situation, the crystal velocity gradient L can be decomposed into a shearing along slip systems and a rigid lattice rotation c
C
α
L = L = SI+ Σγ Ύ"
(1)
a
where Ω is the lattice spin, γ° is the shear rate on slip system a, and Τ is the Schmidt tensor. Here, Τ = b ® η , with b , the slip direction, and n, the slip normal for slip system. In this paper, modeling F S W of stainless steel, which has a face-centered cubic (FCC) structure was investigated. For this material, slip is assumed to occur on the {111} planes and in the <110> directions. As b and η are tied to the crystal lattice, Τ is a function of the crystal orientation, defined here using an angle-axis representation r. The skew part of crystal velocity gradient in Eq. (1) yields an expression for the lattice spin
«=w-xr
(2)
where W is the macroscopic spin and Q = s k e w T . From Ω , the rate of reorientation of the crystal lattice can be written as r = — ω + (ω • r ) r + c o x r
6
·
Materials Processing and Texture
(3)
T e x t u r e E v o l u t i o n d u r i n g Friction Stir W e l d i n g of S t a i n l e s s S t e e l
where ω = v e c t ( f t ) . Equation (2) is the evolution equation for the lattice orientation of a specific crystal. It is integrated for every crystal in a representative volume, or aggregate, over the deformation history that the aggregate experiences as it flows through the Eulerian domain. The collective responses of all the crystals within an aggregate define the texture evolution. The reorientation velocity for a particular crystal at any time depends on the current values of its slip system shearing rates, γ". A viscoplastic constitutive relation is used to relate the shear rate on a slip system to the resolved shear stress
where τ " = Ρ" • a'
τ
(4)
where m is the rate sensitivity, τ is the slip system hardness (assumed identical for all slip systems), γ is a reference strain rate, and τ is the resolved shear stress on slip system obtained as the projection of the crystal stress onto the slip system. P = s y m Τ and a is the deviatoric crystal stress. The rate sensitivity strongly influences the relative activity of the slip systems with the net effect that under monotonie loading, weaker textures result from higher rate sensitivity. For the results shown, a value is given with rate insensitive behavior (TH=0.05). A variable value, ranging from 0.05 at 273 Κ to 0.11 at 1100 Κ was also used to incorporate the effect of greater rate sensitivity with elevation of the temperature. Qualitatively, the textures were similar, but those computed using variable rate sensitivity were weaker, as expected. An expression for the crystal stress follows by combining the symmetric part of crystal velocity gradient equation (1) with Eq. (4): 0
α
c
e =[C r
where the stiffness C
c
, ;
]D''
(5)
depends on the lattice orientation r and is defined as
(6) r
and D is the crystal deviatoric deformation rate. Equation (5) is solved for the crystal stress knowing the current values of the lattice orientation and the crystal deformation rate which under the extended Taylor hypothesis is the same as the macroscopic deformation rate. The slip system shearing rates are subsequently evaluated from Eq. (4) and used in Eq. (2) to compute the lattice spin. 3.2 IDEAL P U R E A N D SIMPLE S H E A R T E X T U R E S During FSW, the basic deformation modes are shearing. Therefore, it is useful to look through textures computed using the extended Taylor assumption that evolve as a consequence of monotonie straining over two ideal deformation modes: simple shear and pure shear. These deformation modes differ in the amount of spin imposed, which is apparent by decomposing the velocity gradient into symmetric and skew parts to give the deformation rate D and the spin W as (7)
L = D +W
The 111 pole figure in Fig. 3(a) shows the pure shear texture. Note that the pure shear texture possesses symmetries about the lines x=y and x=~y, coinciding with the principal axes of the deformation rate. The 111 pole figure in Fig. 3(b) shows the simple shear texture. The texture is rotated
Materials Processing and Texture
·
7
T e x t u r e Evolution d u r i n g Friction Stir W e l d i n g of S t a i n l e s s S t e e l
about the ζ axis from the pure shear texture and the symmetries apparent for pure shear breakdown due to the effect of the superimposed spin. Shear textures and their associated ideal orientations have been studied in several experiments [20, 21] and modeling [22, 2 3 , 2 4 , 2 5 , 26] on cubic polycrystals. Pure shear possesses a number of stable equilibria, as computed using a linearized stability analysis, in the reorientation velocity field. The stable equilibria are listed in Table 2 and are displayed in Fig. 3(c) on a 111 pole figure. Only the stable equilibria, meaning that the divergence of the reorientation velocity field is negative, are shown. Under monotonie loading, the texture will continue to increase in strength at these points, as can be seen by comparison of Fig. 3(a) with Fig. 3(c). The addition of spin to pure shear alters the structure o f the reorientation velocity field. For simple shear, the spin is sufficient to destroy the equilibria. Nevertheless, the texture does strengthen under simple shear in the vicinities of the ideal orientations. The ideal components are listed in Table 3 and shown in Fig. 3(d), along with two ideal fibers, referred to as the A and Β fibers, {111}
and{AW/, respectively. However, as none of the ideal orientations are stable equilibria, the texture will not increase monotonically, but instead will eventually diminish under continued straining. Thus, although the texture may temporarily increase in strength under a simple shear deformation, the strengthening will not persist indefinitely, as discussed in m o r e detail in [25]. As will be shown later, along segments of a streamline, the motion often is reasonably described by this reduced form of the velocity gradient with differing levels of the spin. The strengthening and weakening of the textures can be understood in part by comparison to the ideal shearing textures and by analysis of the flow in terms of level of spin superimposed o n shearing deformations. Such textures might be expected in the F S W process, except that the principal directions of the deformation rate are constantly changing as the material is swept around the tool pin.
Figure 3. For face-centered cubic crystals, 111 pole figures under (a) pure and (b) simple shear modes of deformation and the locations of the (c) stable and (d) ideal texture components for pure shear and simple shear, respectively. Repeated numbers indicate equivalent poles. See Tables 1 and 2 for identification of the poles.
8
·
Materials Processing a n d Texture
T e x t u r e Evolution d u r i n g Friction Stir W e l d i n g of S t a i n l e s s S t e e l
Table II. Coordinates of stable equilibria for pure shear of FCC. Given as Miller indices, Euler angles, and Rodrigues vector. Nairn;
Component
(hkl)[mm]
{r . r i-J
{*,.*. Φ J
x
r
1
C
(0 I Γ)[7 5 5 ]
{45.0. 135.(1. 0 }
{04142. 0.1716. 0.4142}
•>
Bi
(1 1 l ) [ l l 4 15]
{75.0. 125.2. 45.0 }
(-0.3657. 0.3657. 0.0000}
V
D:
(f 1 T ) [ l ! 15 4]
{15.0. 125.2. 3 1 5 . 0 }
{0.3657. -0.3657. 0.0WW}
3
£, E
(1 1 0)[2i 2~1 95]
{72.4. 90.0. 4 5 . 0 }
{0.4142, -0.0642. -0.1551}
(1 1 0)[67 67 30]
{17.6. 90 0 . 4 5 . 0 }
{0.0642.-0.4142. 0.1551}
y
2
3.3 SIMPLIFIED C I R C U L A R S T R E A M L I N E M O D E L (2D) There are dynamic variations of velocity gradients around the tool pin. The simplified circular streamline model provides information on h o w the symmetric deformation rate and the skew spin affect each other. A simplified circular model is described by a flow along a circular path around the tool pin (small window in the upper-left of Fig. 5a). Consider a local coordinate system with base Table III. Coordinates of ideal orientations for simple shear of F C C . Given as Miller indices, Euler angles, and Rodrigues vector. Name
Component 1
-»1
3 4 5 6 1
(ΛΗΗ>·Η·]
{*,.φ.
{',. V
Φυ
r.)
{Ο. 35.26. 45 }
{0.3178 -0.1316 0.4142}
ι ί Γ ΐ ) [ ί ι ο]
{IW, 35.26. 45 }
(I I ί)[2 f 1]
{35 26. 45. Ο }
{-0.3178 0 1316 0.4142} {04142 0 1316 0 317»}
(1 I ί)[Ι
ίο]
(I Γ ! ) [ ϊ Γ I ] .
{144 74. 45. 0 }
{-0.4142 11.1316-CI.3I7X}
ß,
(I I 2")[1 Î 0 ]
{(). 54 74. 45 }
{-0.317)1 0.4142 -0.1316}
β.
(f ί 2 ) [ Γ 1 0]
{ΙΚΟ. 54.74. 45 }
{0 3178 -0.4142 -0.1316}
C
I l 0 01(0 1 I]
{90. 45. 0 }
{0.(ΚΧΚ> -11.4142 0.Ο0ΟΟ}
vectors that are tangent and normal to the circle. The model velocity gradient is constant in this frame. In the global frame, the velocity gradient is found by a change of basis transformation: L = R L
R
(8)
T
where L is a local velocity gradient and R is a rotation tensor relating the local and global coordinate systems. The components of symmetric deformation rate in the global reference frame, D and D , u
xy
exhibit distributions of trigonometric functions. A range of crystal flows developed under arbitrary deformations is generated using different values of the ratio of the macroscopic spin, W , xy
The condition \W I D S)
e(l
| > |/? / r>
to the effective deformation rate, D , rJf
or JFF^ /
.
| corresponds to cases in which the reorientation velocity field has
no stable orientations. The pure and simple shear cases are given in Fig. 4. The texture of pure shear (Fig. 4a) results in a monotonie increase in the intensity of several stable orientations. In this case, \W I D I < \D I D I. In contrast, the texture of simple shear repeatedly strengthens and weakens. The tf
tl]
V
ell
Materials Processing and Texture
·
9
T e x t u r e Evolution d u r i n g Friction Stir W e l d i n g of S t a i n l e s s S t e e l
distance along the streamline is normalized by the length of the streamline, so that the starting and ending points assume normalized distances of 0 to 1, respectively. Texture indices for both pure and simple shear are shown in Fig. 4(c) and 4(d). The texture index measures the sharpness of O D F (orientation distribution function). The texture index of pure shear increases monotonically toward a saturation value. That of simple shear shows regular oscillation between high and low values. 3.4 T E X T U R E PREDICTION USING C R Y S T A L PLASTICITY (2D) The velocity gradients used in the circular streamline model maintain a constant ratio between the magnitudes of the skew and symmetric parts. This results in a regular variation in the texture indices. Unlike the simple circular model, the velocity gradients along the streamlines in FSW modeling vary greatly. In Figure 5(a). the longest streamline, which closely surrounds the tool pin (2D case), can be divided roughly into 12 blocks from LI to L6 and R l to R6. Six of them are located upstream of the tool pin center (from LI to L6) and six are located downstream of the tool pin center (from R 6 to R l ) . Each block corresponds to a region of strengthening texture. The small window in the upper right side of Fig. 5(a) is the enlarged figure around the tool pin. It can be seen that the principal directions of the deformation rate are rotating in a similar manner to those in the simplified circular streamline model. This suggests that the dominant shear plane and shear direction are approximately tangent and normal to the tool pin, respectively. The stagnation point and transition region between blocks also are selected to illustrate the texture evolution. The strengthening of texture around the tool pin is displayed in Fig. 5(b) for each of the 12 blocks in Fig. 5(a).
1
Figure 4. Velocity gradient distributions [ s ] and texture indices around the circle. Local velocity gradients of pure shear and simple shear are shown in (a) and (b), respectively. The corresponding texture indices are shown in (c) and (d). A global reference frame is given by the x-y-z directions, and the center of pole figures corresponds to the global ζ direction. The local frame around the tool pin is given by another set of axes: one axis
10
Materials Processing and Texture
T e x t u r e E v o l u t i o n d u r i n g Friction Stir W e l d i n g of S t a i n l e s s S t e e l
defined by the SPN, one by the S D . and the third by the radial direction (RD). For simplicity, only SPN and SD are shown in Fig. 5(b). The 111 pole figures inside a block show similar textures. Overall, the textures in one block have a simple rotational relationship to the textures of an adjoining block. This relationship is easily understood using the Β fiber. T h e Β fiber, which is located mainly along the SPN direction, rotates according to the tool pin rotation from position to position in Fig. 5(b). For example, the region between blocks L3 and L4 shows that the SPN direction is parallel to the χ direction. Therefore, the Β fiber is located along the χ direction in the pole figure. Fig. 6(a) shows the velocity gradient distribution of the longest streamline around the tool pin in Fig. 5(a). Overall, the distributions are similar to Fig. 4(b), the simple shear case. The condition Ι ^ / ΰ ^ Ι ^ Ι ^ / ΰ ^ Ι refers to a state in which the reorientation velocity field has no stable orientations. The oscillations of the texture index in Fig. 6(b) correspond to strengthening, weakening, and restrengthening of the texture. A two-dimensional model is based on a plane strain deformation, and is limited to reflect the more complicated deformation patterns due to the influence of the pin thread and shoulder. 3.5 T E X T U R E PREDICTIONS USING C R Y S T A L PLASTICITY (3D) For a simplified two-dimensional, plane strain model, the out-of-plane strain components are identically zero. Therefore, the texture of F S W is expected to have a mixture of plane strain (or pure shear) and simple shear components. In a three-dimensional model, more complex textures are possible. Fig. 7 shows the textures predicted by a three-dimensional simulation of FSW. Each number specifies a position in the outlet where the textures are calculated. Position 2 is located at the end of one of the longest streamlines, passing areas affected by both the tool shoulder and pin. Such material points start from the top surface, spiral down under the tool shoulder and around the tool pin, and finally leave the stirred zone in the bottom part of the plate.
Figure 5. T h e longest streamline and the position for the 111 pole figures. Simplified circular streamline model resembles the deformation around the tool pin. Each symbol represents the regions from LI to L6 and R6 to R l . (a) overall model domain, (b) 111 pole figures for the regions of L and R.
Materials Processing a n d Texture
11
T e x t u r e Evolution d u r i n g Friction Stir W e l d i n g of S t a i n l e s s S t e e l
Figure 6. Velocity gradients and texture indices around the tool pin. Symmetric and skew parts of velocity gradient are shown, (a) velocity gradient and (b) texture index. Those material points experience severe deformation, traveling a longer distance and spending a longer time in the stirred zone than other material points. Fig. 8 shows the velocity gradient and texture index for the material point associated with Position 2 in Fig. 7. The streamline ending at Position 2 is split into two parts, the region under the shoulder and the region around the pin. The material point makes several loops around the tool pin. Under the shoulder, the magnitudes of spin and deformation rate are similar. Those distributions create minor oscillations in the texture index. Around the pin, the magnitudes of deformation rate are larger than those of spin, (\W I D 1 < I D |; / 1 - j ). The texture index shows m
cl/
tll
higher and severer oscillation. In the transition zone between tool shoulder and pin (normalized distance of 0.6 - 0.8), the magnitude of spin is higher than that of the deformation rate. This results in a texture index reduction.
12
·
Materials Processing and Texture
T e x t u r e E v o l u t i o n d u r i n g Friction Stir W e l d i n g of S t a i n l e s s S t e e l
shoulder β
7
^4;
10
plate
pin 1
2
=3=
s
Figure 7. Pole figures predicted from F S W simulations.
'S
Normalized distance
1
Figure 8. Velocity gradient distributions ( s ) and texture indices around the tool pin and tool shoulder.
3.6 E X P E R I M E N T S (EBSD) The results of EBSD measurements are shown in Fig. 9. Measured 111 pole figures show two features of texture during FSW: the shear texture component and its rotation with position. Figs. 9(d), (e), (f), and (g) show pole figures with single orientations and backgrounds. The single orientation in the
Materials Processing and Texture
·
13
T e x t u r e Evolution d u r i n g Friction Stir W e l d i n g of S t a i n l e s s S t e e l
pole figures appears to rotate as the position on the specimen varies. Major peaks in the pole figure can be understood as the rotation of an ideally-stable shear orientation, C. These major peaks, or rotated stable shear orientations, were also found in the texture prediction using crystal plasticity during FSW. In Fig. 7, Positions 1,2,4, and 5 show a single orientation distribution and a fiber texture. In particular, both Position 2 and 4 have the longest streamlines, traversing both the tool shoulder and the tool pin.
Figure 9. 111 pole figures measured from the cross section of friction stirred plates.
4.
CONCLUSION
Texture evolution during F S W was investigated using the simple circular streamline model, two- and three-dimensional Eulerian finite element models, and E B S D measurements. 1. Texture evolution is predicted from velocity gradient histories along streamlines. In the plane strain case, the interplay between the magnitudes of deformation rate and spin determines whether the texture will stabilize. The F E models indicate that the texture forms, breaks up, and reforms repeatedly and rapidly. 2. In three-dimensional modeling, both the tool pin and shoulder are considered. The axial component of the pin traction and the tangential component of shoulder traction both act to produce three dimensional flow. A pin thread angle of 5 deg is sufficient to establish substantial motion of the thickness direction. 3. The material points that experience the most severe deformations ultimately display a texture that is dominated by a single orientation. These points show a gradual increase in the texture index although the texture repeated weakens and strengthens. Other points experience less severe deformations and display a texture described best as a fiber distribution. 4. The ratio of the magnitude of the spin to the magnitude of the deformation rate is central to the texture stability. If this ratio is sufficiently high, stable texture components do not exist and the texture evolution is characterized by periods of strengthening and weakening. 5. Weld and base zone microstructures, as measured using E B S D , display differences in grain size, grain morphology, and distributions of annealing twins. Material from the stirred zones displays both single-component and fiber textures consistent with the simulated textures.
14
Materials Processing and Texture
T e x t u r e E v o l u t i o n d u r i n g Friction Stir W e l d i n g of S t a i n l e s s S t e e l
ACKNOWLEDGEMENT The author would like to thank D.E. Boyce and Prof. P.R. Dawson of Cornell University. REFERENCES [1] Boyce, D. E., Dawson, P. R., Sidle, B., and Gnaupel-Herold, T., A Multiscale Methodology for Deformation Modeling. Comput. Mater. Sei., 3 8 , pp. 158-175 (2006). [2] Field, D., Nelson, T., Hovanski, Y., and Jala, K., Heterogeneity of Crystallographic Texture in Friction Stir Welds of Aluminum, Metall. Mater. Trans. A, 32A(11), pp. 2 8 6 9 - 2 8 7 7 (2001). [3] Sato, Y., Kokawa, H., Ikeda, K., Enomoto, M., Jogan, S., and Hashimoto, T., Microtexture in the Friction-Stir Weld of an Aluminum Alloy. Metall. Mater. Trans. A, 32A(4), pp. 9 4 1 - 9 4 8 (2001). [4] Fonda, R., Bingert, J., and Colligan, K., Development of Grain Structure During Friction Stir Welding, Scr. Mater., 5 1 , p p . 2 4 3 - 2 4 8 (2004). [5] Prangnell, P. B., and Heason. C. P., Grain Structure Formation During Friction Stir Welding Observed by the 'Stop Action Technique', Acta Mater., 53, pp. 3179-3192 (2005). [6] Sato, Y. S., Nelson, T., and Sterling, C. J.. 2005, Recrystallization in Type 304L Stainless Steel During Friction Stirring, Acta Mater., 5 3 , pp. 6 3 7 - 6 4 5 . [7]. K. Wierzbanowski and A. Clement, Rotation Field and Continuity Equation for Texture Evolution, Cryst. Res. Technol., 1984, vol. 19, pp. 201-12. [8]. A. Kumar and P.R. Dawson, The Simulation of Texture Evolution with Finite Elements over Orientation Space - I. Development. Comput. Methods Appl. Meek Eng.. 1996, vol. 130, pp. 247-61. [9], L. Torn, P. Gilormini, and J. Jonas, Effect of rate sensitivity on the stability of torsion textures, Acta Metall., 1988, vol. 36, pp. 3077-91. [10]. L. Toth, K. Neale, and J. Jonas. Stress response and persistence characteristics of the ideal orientations of shear texture. Acta Metall., 1989. vol. 37, pp. 2197-210. [11]. A. Kumar: Ph.D. Thesis. Polycrystal Modeling with Finite Elements over Orientation Space, Cornell University, Ithaca, N Y . 1996. [12] Cho, J.-H., Boyce. D., and Dawson, P. R., 2005, Modeling Strain Hardening and Texture Evolution in Friction Stir Welding of Stainless Steel, Mater. Sei. Eng., A. 398, pp. 1 4 6 - 1 6 3 . [13] Cho, J.-H., Boyce, D. E., and Dawson, P. R., 2007, Modeling Strain Hardening During Friction Stir Welding of Stainless Steel, Modell. Simul. Mater. Sei. Eng., 15, pp. 4 8 6 - 4 8 9 . [14] Li, Y.. Murr, L., and McClure, J., 1999, Flow Visualization and Residual Microstructures Associated With the Friction-Stir Welding, Mater. Sei. Eng., A, A 2 7 1 , pp. 2 1 3 - 2 2 3 . [15] Guerra, M., Schmidt, C . McClure, J., Murr, L., and Nunes, Α.. 2003, Flow Patterns During Friction Stir Welding, Mater. Charact., 49, pp. 9 5 - 1 0 1 . [16] Mathur, K., and Dawson, P., 1989, O n Modeling the Development of Crystallographic Texture in Bulk Forming Processes, Int. J. Plast., 5, pp. 6 7 - 9 4 . [17] Dawson. P. R., and Wenk, H.-R., 2000. Texturing the Upper Mantle During Convection. Philos. Mag. A, 80(3), pp. 5 7 3 - 5 9 8 . [18] Kocks, U. F., Tome, C. N., and Wenk, H.-R., 1998, Texture and Anisotropy, Cambridge University Press, Cambridge. [19] Taylor, G., 1938, Plastic Strains in Metals, J. Inst. Met., 62, pp. 3 0 7 - 3 2 4 . [20] Sevillano, J. G., Houtte, P. V., and Aernoudt, E., 1977, The Contribution of Macroscopic Shear Bands to the Rolling Texture of F C C Metals, Scr. Metall., 1 1 , pp. 5 8 1 - 5 8 5 . [21] Montheillet, F., Gilormini, P., and Jonas, J., 1984, Relation Between Axial Stresses and Texture Development During Torsion Testing: A Simplified Theory, Acta Metall., 3 3 , pp. 7 0 5 - 7 1 7 .
Materials Processing and Texture
15
T e x t u r e Evolution d u r i n g Friction Stir W e l d i n g of S t a i n l e s s S t e e l
[22] Toth, L., Gilormini, P., and Jonas, J., 1988, Effect of Rate Sensitivity on the Stability of Torsion Textures, Acta Metall., 3 6 , pp. 3 0 7 7 - 3 0 9 1 . [23] Toth, L., Neale, K., and Jonas, J., 1989, Stress Response and Persistence Characteristics of the Ideal Orientations of Shear Texture, Acta Metall., 3 7 _ 8 _ , pp. 2 1 9 7 - 2 2 1 0 . [24] Canova, G., Kocks, U., and Jonas, J., 1984, Theory of Torsion Texture Development, Acta Metall., 3 2 , pp. 2 1 1 - 2 2 6 . [25] Kumar, Α., 1996, Polycrystal Modeling With Finite Elements Over Orientation Space, Ph.D. thesis, Cornell University, N e w York. [26] Kumar, Α., and Dawson, P., 2000, Computational Modeling of F.C.C Deformation Textures Over Rodrigues' Space, Acta Mater., 48, pp. 2 7 1 9 - 2 7 3 6 .
16
Materials Processing a n d Texture
T E X T U R E D E V E L O P M E N T IN A L U M I N U M F R I C T I O N STIR W E L D S Richard W . F o n d a and Keith Ε. Knipling US Naval Research Laboratory Code 6 3 5 6 Washington, D C , 2 0 3 7 5 John F. Bingert Los A l a m o s National Laboratory M S G 7 5 5 , P.O. Box 1663 Los A l a m o s , N M , 87545 Anthony P. R e y n o l d s and Wei T a n g University of South Carolina College of Engineering Columbia, S C , 2 9 2 0 8 Kevin J. Colligan Concurrent T e c h n o l o g i e s Corporation 27980 K i m Drive Harvest, A L 3 5 7 4 9 John A. Wert Rise National Laboratory Frederiksborgvej 3 9 9 4000 Roskilde, D e n m a r k ABSTRACT This study reveals the d e v e l o p m e n t of texture and microstructure that occurs during friction stir welding of t w o a l u m i n u m alloys. Friction stir w e l d s of single crystal a l u m i n u m and 2195 aluminum w e r e investigated using a stop-action technique, w h e r e the microstructure surrounding the tool w a s preserved by immediately q u e n c h i n g the weld end u p o n completion. Plan view sections of these w e l d terminations reveal that slight variations in the crystallographic orientation within the base material h a v e different susceptibilities to the shear deformation produced by the rotating tool. Similarly, the misorientations generated during subgrain development can also induce variations in the lattice rotation rate. The divergent orientations that develop can have different susceptibilities to both rotation and grain subdivision. Lattice rotations continue until achieving a readily-sheared orientation, after w h i c h n o further rotation is observed. The final texture that develops near the tool corresponds to the Β and Β c o m p o n e n t s of the simple shear texture. In the 2 1 9 5 alloy, the deposited weld also exhibits a simple shear texture, but it varies with a periodicity of the tool a d v a n c e per revolution b e t w e e n a pure BIB c o m p o n e n t and a mixture of the BIB and C c o m p o n e n t s . T h u s , the d e v e l o p m e n t of textures ahead of the tool can be correlated directly to the increasing temperature and shear deformation of those regions, while the periodic texture variations of the deposited w e l d indicate a systematically varying c o m p o n e n t in the friction stir welding process. INTRODUCTION Friction stir welding ( F S W ) is a solid-state j o i n i n g technique that w a s developed in 1991 ' and has rapidly b e c o m e a viable process for j o i n i n g a l u m i n u m alloys. D u r i n g the F S W process,
17
T e x t u r e D e v e l o p m e n t in A l u m i n u m Friction Stir W e l d s
a rotating tool is p l u n g e d into the j o i n t line, where it heats and plastically deforms ( ' s t i r s ' ) t h e surrounding material. A s the rotating tool is traversed along the joint, the material from both sides are stirred together and t h e n deposited in the w a k e of the tool, thereby forming the weld. Previous studies h a v e e x a m i n e d t h e grain structures and textures that are generated in t h e deposited weld during F S W of polycrystalline a l u m i n u m . " T h e c o m p l e x thermal and strain gradients near the tool, coupled with the material transport around the tool, h o w e v e r , severely complicates the forensic determination of the grain and textural evolution that transformed the original base plate material to the deposited weld. W e have e m p l o y e d stop-action weld techniques, during w h i c h the welding process is halted and the weld is immediately quenched, to preserve the microstructure present around t h e tool during the F S W process. This technique allows a direct study of the deformation m e c h a n i s m s and microstructural evolution that occur during F S W . " " T h e current p a p e r will expand o n some of these s t u d i e s ' ' ' w i t h a focus o n the textural evolution that o c c u r s a h e a d of the tool during F S W of a l u m i n u m alloys. 2
10
1 8
3
4
8
EXPERIMENTAL Friction stir w e l d s w e r e fabricated in a bead-on-plate (without a w e l d j o i n t line) configuration in t w o alloys. T h e first alloy, a high-purity a l u m i n u m single crystal, permits study of the initial influence o f the rotating tool o n a single crystallographic variant of the base material. T h e second alloy, a polycrystalline 2195 a l u m i n u m alloy, is m u c h m o r e resistant to recrystallization and is therefore m o r e suitable to e x a m i n i n g the evolution that occurs closer to the tool. Fabrication and analysis of these samples is discussed b e l o w . Single Crystal A l u m i n u m Friction Stir W e l d T h e single crystal samples w e r e taken from a 20 χ 2 0 χ 100 m m crystal g r o w n from p u r e a l u m i n u m with an arbitrary crystal orientation. Four 4 m m thick slabs were cut from this single crystal and inserted into rectangular holes machined in the center of 4 χ 120 χ 2 4 0 m m sheets of annealed a l u m i n u m A A 1 0 5 0 . T h e single crystal slabs w e r e held in p l a c e b y fusion w e l d s at each corner. Friction stir welding w a s performed with a s m o o t h cylindrical tool with a convex, radiused termination tilted to a 2.5° lead angle in Z-position control at a welding speed of 50 m m min"' (2 in min"') and a tool rotation rate of 8 0 0 rpm. T h e tool had a shoulder diameter of 12.7 m m , a probe diameter of 5.1 m m , and a probe length of 3.2 m m . W e l d i n g w a s started in the polycrystalline A A 1 0 5 0 material and terminated at the centerline of the single crystal to m i n i m i z e effects from the single crystal-to-AA1050 boundary. T h e weld w a s immediately quenched with cold water after withdrawal of the tool to m i n i m i z e subsequent microstructural evolution. Similarly, to preserve the microstructure surrounding the first of the t w o w e l d s m a d e in each panel, that w e l d w a s draped with a water-soaked cloth during fabrication of the second weld. 3
3
Four w e l d s w e r e investigated in order to evaluate the effect o f crystal orientation o n the microstructural and textural evolution during F S W . T h e orientations of these welds are illustrated schematically
18
F i g u r e 1. Schematic s h o w i n g the four welding orientations of the single crystal a l u m i n u m friction stir welds.
Materials Processing and Texture
T e x t u r e D e v e l o p m e n t in A l u m i n u m Friction Stir W e l d s
in Figure 1, although the actual w e l d s w e r e m a d e o n three separate panels. The crystallographic orientation of the single crystal relative to the panel axes, P , P , and P , and the welding orientations (at the front of the tool and assuming a cylindrical shear around the tool) is provided in Table I. x
y
z
Table I: Crystal a x e s parallel to the panel coordinates and to the w e l d axes. Panel Axes Crystal axis
Px [0.83, -0.46, -0.31]
Pv [0.18, 0.76, -0.63]
Pz [0.52, 0.47, 0.72]
Weld Axes Weld 1 Weld 2 Weld 3 Weld 4
Tool advance [-0.83,0.46, 0.311 [0.83, -0.46, -0.311 [-0.18,-0.76, 0.631 [0.46, 0.86, -0.231
Transverse [-0.18, -0.76, 0.63] [0.18, 0.76, -0.63] [ 0 . 8 3 , - 0 . 4 6 , -0.311 [-0.72, - 0 . 2 1 , 0 . 6 6 ]
Rotation Axis [0.52, 0.47, 0.721 [0.52, 0.47, 0.721 [0.52, 0.47,0.721 [0.52, 0.47, 0.721
R e g i o n s immediately ahead of the welding tool w e r e examined at the midplane thickness in each weld. S p e c i m e n s w e r e prepared by sectioning the exit hole of each w e l d (where the welding tool w a s extracted at the end of the w e l d ) and then thinning that sample to half the plate thickness. Electron backscattered diffraction ( E B S D ) scans w e r e performed with a 2 χ 2 μ π ι grid spacing in a region extending 2.0 m m from the tool exit hole along t h e welding direction in the plan-view orientation. A n additional scan w a s performed 6-8 m m a h e a d of the tool in each weld to confirm the relative base plate orientations. 2
Polycrystalline A l u m i n u m Friction Stir Weld The polycrystalline friction stir w e l d w a s prepared as a bead-on-plate weld in a single 25 mm-thick plate of the aluminum-lithium alloy 2 1 9 5 . T h i s alloy is ideally suited for continuously tracking the grain structure and crystallographic texture across the weld-affected regions around the tool d u e to its strong initial texture and resistance to recrystallization. W e l d i n g was performed under load control at a rotational speed of 180 r p m and a translational speed of 1.7 m m s"' (4 in min"') using a tool with a flat, scrolled shoulder and a threaded, truncated-cone pin with three flats. A s w i t h the single crystal friction stir weld, this w e l d w a s immediately quenched u p o n completion of the weld to preserve the microstructure around the tool. Because of its re-entrant features, however, the tool w a s not extracted prior to quenching. T h e end of the weld (containing the e m b e d d e d pin) w a s sliced parallel to the top surface at a depth of 10.4 m m b e l o w the surface by electro-discharge machining. T h i s p l a n - v i e w surface w a s mechanically and then electrochemically (30 s at - 2 0 °C and 15V in 3 0 % H N 0 in methanol) polished for examination by E B S D in a L E O 1550 S E M . 3
RESULTS AND DISCUSSION Friction Stir W e l d i n g of Single Crystal A l u m i n u m Since the single crystal specimen contains a single crystallographic orientation with no grain boundaries, any observed textural evolution or grain boundary development is directly attributable to the friction stir welding process. A s indicated in Table I, the initial base plate w a s oriented with the plate n o r m a l close to [111]. E B S D scans of the single crystal base plate (Figure 2) reveal s o m e slight m o d u l a t i o n s in the base plate orientation that are generally aligned along the short dimension of the rectangular single crystal. T h e s e misorientations will be s h o w n
Materials Processing and Texture
19
T e x t u r e D e v e l o p m e n t in A l u m i n u m Friction Stir W e l d s
to develop into larger misorientations and eventually form n e w grains as the single crystal b e c o m e s influenced by the heat and deformation surrounding the F S W tool. Figure 3 is a p l a n - v i e w E B S D scan from ahead of the tool in W e l d 3 (see Figure 1). T h e action of t h e rotating tool causes the pre-existing orientation m o d u l a t i o n s in the base plate t o b e c o m e m o r e pronounced, . , . , . , , Figure 2 . Crystal orientation m a p and pole . , , ,, .1 . . increasing in both m a g n i t u d e and frequency. ° . , . : figures o f the single crystal base plate material. This is evident m Figure 3 as an increase in contrast and density of those orientation m o d u l a t i o n s from the right side of the image (far from the tool) t o w a r d s the left (closer to the tool). Pole figures from these regions (Figures 4 and 5 from the regions indicated in Figure 3) reveal additional details o f this evolution. T h e base plate has a distinct crystal orientation with little orientation spread (Figure 4), as expected for a single crystal. Closer to the tool (see Figure 5), the material generally experiences a slight rotation in response to the deformation field surrounding the rotating tool. At the s a m e time, h o w e v e r , s o m e orientations experience significant rotations from t h e initial base plate orientation. T h e s e rotated orientations correspond t o the evolving orientation modulations discussed a b o v e , w h i c h increase in both misorientation and n u m b e r nearer to the tool. Significant misorientations d e v e l o p b e t w e e n these regions and the base plate, eventually p r o d u c i n g the high-angle grain boundaries visible approximately 1.5 m m from the tool (near " 5 " in Figure 3). Analysis o f the other weld orientations (see F i g u r e 1), demonstrate that the grain boundaries that d e v e l o p from the single crystal maintain a consistent orientation despite c h a n g e s in the weld orientations. T h u s , these grain b o u n d a r i e s e v o l v e from the small orientation variations present in the original single crystal rather than in r e s p o n s e t o a particular deformation orientation. T h e increased variation in orientation that develops with increasing deformation is consistent with shear deformation m o d e l s ' " , w h i c h predict that lattice rotation rates vary significantly as a function of crystallographic orientation. Slight differences in orientation c a n thus lead to different rotation rates between those regions, further accentuating any orientation difference. In accord with these m o d e l s , the material ahead of the tool in these single crystal w e l d s generally rotates slowly in r e s p o n s e to the stress (Figures 4-6) while the slightly misoriented regions rotate at a faster rate to generate a spread of orientations misoriented by u p to 50° from the base plate orientation (see Figure 5). c
J
D
9
J
21
Figure 3 . E B S D crystal direction m a p of the developing microstructure a h e a d of the tool in W e l d 3 (see Figure 1) s h o w i n g the locations used for Figures 4, 5, 6, and 8. W e l d i n g direction is to right and exit hole is s h o w n at left.
20
·
Materials Processing and Texture
T e x t u r e D e v e l o p m e n t in A l u m i n u m Friction Stir W e l d s
A p p r o x i m a t e l y 1 m m from the tool, large elongated grains are o b s e r v e d aligned along the predominant stress field. T h e s e grains a p p e a r to have a r a n d o m orientation and have much lower strain t h a n t h e surrounding material, as indicated by their higher E B S D pattern quality. They appear to result from conventional, static recrystallization , during w h i c h the heat and deformation introduced during the F S W process facilitate the nucleation of n e w grains, which then g r o w into t h e surrounding, deformed microstructure. While the statically recrystallized grains in t h e outer regions of t h e weld-affected zone contain few, if a n y , internal boundaries, the 25
Figure 4. 101 and 111 pole figures obtained from the region indicated by a " 4 " in Figure 3 , s h o w i n g the undeformed texture present far from the friction stir welding tool.
Figure 5. 101 and 111 pole figures obtained from the region indicated by a " 5 " in Figure 3 , s h o w i n g the initial evolution of texture from the base plate orientation in response to the thermal and deformation fields s u r r o u n d i n g t h e friction stir w e l d i n g tool.
Figure 6. 101 and 111 pole figures obtained from the region indicated by a " 6 " in Figure 3 (the fine grains b e t w e e n the large, statically recrystallized grains), showing d e v e l o p m e n t of a terminal orientation after rotation from the b a s e plate orientation.
n u m b e r of low- and m e d i u m - a n g l e grain b o u n d a r i e s increases in these grains as they approach the welding tool. This indicates that static recrystallization occurs during the initial stages of deformation, b u t these recrystallized grains b e c o m e deformed as they approach the tool and are influenced by the increasing rotational deformation field. C l o s e to the tool, no evidence of these statically recrystallized grains remains, presumably because the continued subdivision of these grains m a k e s them indistinguishable from the rest of the dynamically recrystallized microstructure. T h e extent o f the weld-affected region (and the v o l u m e that is consumed by the statically recrystallized grains) varies as a function of weld orientation. W h i l e différences in weld orientation within the single crystal are expected to p r o d u c e microstructural variations ahead of the tool d u e to differences in susceptibility to deformation, recovery- of the deformation structures, and recrystallization o f those structures, other factors m a y also influence the observed structure. In particular, the shape of the single crystal slab within the weld panel accounts for m u c h of this variation. T h e reduced thermal conductivity across the A A 1 0 5 0 plate/single crystal interface limits heating of the single crystal until the F S W tool crosses that interface. T h u s , i j n t h e single between Welds 2, 3 , ^ environment at the end of t h e w e l d
m
c
e
r
a
n
i
n
d i s t a n c e
y
s
t
a
l
u
e
d
f
L
w
4
h
i
c
( s e e
n
c
e
o
n
w
h
e
d
e
d
w i t n
v a r i e s
F
m
i
e
g
u
r
e
h
a
s
a
s t r o n g
t h e r r n a l
Materials Processing and Texture
·
21
T e x t u r e D e v e l o p m e n t in A l u m i n u m Friction Stir W e l d s
T h e grain and texture evolution from the original single crystal base plate can continue to be tracked b e t w e e n the statically recrystallized grains. A t the outer edge of this region, the b a n d s of small grains that originated from the base plate orientation m o d u l a t i o n s b e c o m e deflected by the rotational deformation field and rotate towards the circumferential flow direction of the stir zone (see Figure 3). T h e evolution of these grains continues t o w a r d s the tool between the large, statically recrystallized grains. Surprisingly, the grains in this region do not exhibit a continued rotation from the base plate orientation, as m i g h t be expected from Figure 5. Instead, they rotate until a specific terminal orientation is reached (see Figure 6). While m o d e l s ' " indicate that no orientation is stable during torsion deformation, they also predict that lattice rotation rates depend strongly o n orientation and that rotating grains near particular ideal orientations can remain near those orientations o v e r appreciable strains. Indeed, w h e n the pole figures of c o m p a r a b l e regions in Welds 2, 3 , and 4 are c o m p a r e d in the shear frame o f reference (Figure 7), t w o of these w e l d s ( W e l d s 2 and 4 ) exhibit terminal orientations w h i c h align < 1 1 0 > along the shear direction and/or {111} with the shear plane n o r m a l . These t w o orientations correspond respectively to the and '-4' partial fibers of an ideal shear t e x t u r e and are predicted to have severely retarded rotations rates. Closer to the tool, w h e r e the statically-recrystallized grains begin t o subdivide from the heat and deformation of w e l d i n g , the texture is dominated by the ' Β ' and ' Β ' partial fibers of the (Figure 8). This ideal shear t e x t u r e texture, characterized by a < 1 1 0 > shear direction, a {112} shear plane normal, 9
21
F i g u r e 7. 101 and 111 pole figures from the (a) W e l d 2, ( b ) W e l d 3 , a n d (c) Weld 4 orientations at locations similar to that in Figure 6. T h e s e pole figures h a v e b e e n rotated to orient the center of rotation at the pole figure center. Shear direction and shear plane normal orientations are approximated from microstructural features in the corresponding E B S D scans. Arrows show the rotations experienced by the developing (sub)grains.
23
23
22
Materials Processing and Texture
Figure 8. 101 and 111 pole figures from grainrefined region ahead of the tool in W e l d 1, exhibiting the ' Β ' and ' Β ' partial fibers of the ideal shear texture characteristic of high-strain, high-temperature shear deformation.
T e x t u r e D e v e l o p m e n t in A l u m i n u m Friction Stir W e l d s
24
25
and a < 1 1 1 > radial direction, h a s also been o b s e r v e d in high-temperature t o r s i o n , ' equal channel angular e x t r u s i o n , and other friction stir w e l d s . ' ' This is the d o m i n a n t texture throughout most of the grain-refined region surrounding the tool. The shear frame of reference is similar to that o b s e r v e d in t h e outer portions of the weld, indicating a consistent orientation of the shear deformation field irrespective of distance from the tool. 2 6 2 7
7
10
1 4
A scan of the initial deformation occurring a h e a d of t h e tool in Weld 2 (Figure 9) revealed an interesting feature. A s in t h e other weld orientations, the b a s e plate and slightly misoriented regions rotate in accord with the weld-induced deformation, albeit at different rotation rates. A misorientation thus develops b e t w e e n the p r e d o m i n a n t single crystal orientation and these Figure 9. E B S D crystal direction m a p and 111 pole misoriented regions (aligned figures from ahead of the tool in W e l d 2 showing the vertically in this w e l d orientation). p r e d o m i n a n t counterclockwise evolution towards a The rotation continues until achieving terminal orientation (black arrows) as well as localized a terminal orientation, s h o w n in formation o f deformation b a n d s and their lattice black, w h i c h aligns {111} with the rotation in the opposite (clockwise) direction. shear p l a n e n o r m a l as w a s s h o w n in Clockwise-rotating tool is located to left o f Figure. Figure 7a. H o w e v e r , s o m e localized deformation b a n d s (the light bands in Figure 9) experience a rotation counter to the predominant rotation. T h e orientation in these b a n d s evolves from the initial single crystal orientation in a clockwise direction until reaching an orientation with < 1 0 1 > aligned along the shear direction (See Figure 7a). This terminal orientation a p p e a r s in the pole figures and within these deformation b a n d s as a d a r k grey color. Friction Stir W e l d i n g of Polycrystalline 2195 A l u m i n u m T h e 2 1 9 5 alloy is m u c h m o r e resistant to recrystallization and is therefore m o r e suitable for studying t h e microstructural and textural evolutions that occur in the refined grains closer to the tool during F S W . Figure 10, an E B S D scan from the a d v a n c i n g side of the tool, demonstrates that variations in the initial grain orientation and gradients in shear stresses (and temperature)
Figure 10. E B S D crystal direction m a p showing the d e v e l o p m e n t of subgrains and grains from the 2 1 9 5 base plate (left) t o w a r d s t h e clockwise-rotating tool (to right of Figure), and plot of the local and cumulative misorientations that d e v e l o p a l o n g the indicated subgrain.
Materials Processing and Texture
23
T e x t u r e D e v e l o p m e n t in A l u m i n u m Friction Stir W e l d s
w i t h proximity t o t h e tool c a u s e the large grain at the left to subdivide t o w a r d s t h e tool (towards the right). T h i s differs from the textural evolution observed in the single crystal friction stir weld, where initial misorientations in the base plate gave rise to different rotation rates and t h u s divergent orientations. In this case, deformation-induced subgrain formation is active, w h i c h m a y result from t h e smaller initial orientation spread within the grains o f the polycrystalline 2195 alloy t h a n in the single crystal. T h e resultant subgrains b e c o m e elongated a l o n g t h e macroscopic shear direction and d e v e l o p into elongated, parallel subgrains (arrowed). A trace along an individual subgrain (highlighted) reveals that this subgrain u n d e r g o e s m o r e than 30° of rotation across a distance of 6 0 0 μπι. T h i s m a c r o s c o p i c lattice rotation evolves in a continuous manner and is accomplished solely by small angle misorientations (<5°). Other subgrains in this image experience a similar evolution with s o m e subgrains d e v e l o p i n g m o r e t h a n 50° of internal misorientations. T h e s e elongated subgrains are separated by boundaries (arrowed) w i t h misorientations that can range to m o r e than 15°. The grain and texture evolution occurring ahead of the tool (Figure 11 ) are similar to that observed in Figure 10, although they are m o r e c o m p r e s s e d due to the steeper deformation and thermal gradients at that location. Similar subgrain and grain d e v e l o p m e n t s to those discussed a b o v e are evident t o w a r d s the left side o f this image. Even closer to the tool, b a n d s of alternating refined grains and large, undeformed grains appear to evolve. T h e larger grains h a v e an orientation that aligns < 1 0 1 > along the vertical shear direction while the refined grains h a v e a m o r e r a n d o m orientation. T h u s , these periodic b a n d s of refined grains appear to arise from regions (possibly parts of t h e s a m e grain that developed a different orientation due t o variations in the ,, , rotation rate or else regions from a F
j
g
u
r
e
E
ß
S
D
c r y s t a
d
i
r
e
a
i
o
n
m
a
p
f
r
o
m
a
h
e
a
d
o
f
separate gram that w a s elongated ^ „üciostnictunil i i o n from the by the shear deformation) that are , , _ oriented in a m o r e deformationsusceptible orientation than the adjacent large grains. Small subgrains d e v e l o p within the elongated subgrains w h i c h t h e n gradually and continually evolve into the fine, e q u i a x e d grains near t h e tool. Initially, small, e q u i a x e d subgrains d e v e l o p with a slight misorientation from the surrounding material (boundary misorientations o f 2-5°), and then increase further in misorientation as they approach the tool. Throughout their d e v e l o p m e n t , the subgrains r e m a i n similar in both size and shape to the m o r e pronounced subgrains and grains observed closer to the tool. T h e development of these t
2
24
·
h
e
1
s
9
5
Materials Processing and Texture
b
a
s
e
h
p
o
w
i
a t e
m
fleft)
e
e v o
t £ )
t h e
t 0 0
( r f g h t )
u t
T e x t u r e D e v e l o p m e n t in A l u m i n u m Friction Stir W e l d s
subgrains thus appears to o c c u r through a continuous d y n a m i c recrystallization ( C D R X ) m e c h a n i s m , consisting of a continuous introduction o f deformation into the structure coupled with rapid recovery of that strain to the grain/subgrain walls, and migration o f the grain/subgrain boundaries to maintain the e q u i a x e d shape. Continued deformation and recovery gradually increases the misorientation of those boundaries to generate the fine, equiaxed grain structure near the tool t h r o u g h the same p r o c e s s . T h i s similarity suggests a single grain subdivision m e c h a n i s m operating across all the regions a r o u n d the tool. 9
28
Figure 12. 110 and 111 pole figures from eight locations around the front o f t h e welding tool, showing the textural evolution that occurs in the grain-refined material adjacent to the tool a s it is transported around the tool. The reference frame of the pole figures is rotated to m a i n t a i n the horizontal axis aligned with the tangential direction o f the tool (as is the case directly ahead of the tool), roughly aligning the p r e s u m e d shear direction with the horizontal axis.
A texture consisting of the S and Β c o m p o n e n t s o f the simple shear texture is again observed in the fine-grained material ahead of the tool (along the right side of Figure 11), similar to the observations in the single crystal friction stir w e l d s . This texture is indicative of very large shear strains, suggesting that a mature texture develops prior to the material being swept around the tool and deposited in its w a k e . However, analysis of eight separate regions around the leading edge of the tool (Figure 12) reveals the BIB texture in only s o m e o f the regions adjacent to t h e tool. T h e strength of this
texture generally increases as the material is swept a r o u n d the tool from the advancing side (left) to the retreating side (right). This parallels the accumulation o f deformation that is expected to and accompany this material transport. T h e s e results differ from t h e results of m o d e l s e x p e r i m e n t s on the texture evolution around the tool in austenitic stainless steel F S W s , which reveal that the shear texture is repeatedly formed and broken up as t h e material is transported around the F S W tool. Possible reasons for this discrepancy include the m u c h larger tool used in the current study and/or differences in the intrinsic properties (resistance to recrystallization, thermal conductivity, etc.) of the t w o base plate materials. 29
30
T h e β and Β c o m p o n e n t s r e m a i n d o m i n a n t c o m p o n e n t s of the observed texture in the ' (Figure 13) similar to deposited w e l d behind the tool. H o w e v e r , there is a textural b a n d i n g the "Onion ring" microstructural banding often o b s e r v e d in friction stir welds. The texture varies with a periodicity of the tool a d v a n c e per revolution b e t w e e n a pure BIB c o m p o n e n t and a mixture o f the BIB and C c o m p o n e n t s of the simple shear t e x t u r e . T h e presence of the C c o m p o n e n t in the deposited weld indicates that these regions were subjected t o a smaller, but still substantial, integrated strain. T h u s , the b a n d s containing the C c o m p o n e n t either h a v e u n d e r g o n e less total deformation or else arose through deformation of material that had undergone recrystallization t h r o u g h a p r o c e s s that generates n e w grain orientations. 1 5 1 6
3 1
31
Materials Processing a n d Texture
·
25
T e x t u r e D e v e l o p m e n t in A l u m i n u m Friction Stir W e l d s
Figure 13: E B S D crystal direction m a p from behind the tool, s h o w i n g the periodic textural variation that d e v e l o p s in the deposited w e l d and the typical pole figures and orientation distribution functions from each o f the alternating textures. O D F contours at 0.607, 1.000,1.648, 2.714, 4.472, 7.368, 12.139, A n d 20.000. SUMMARY A n a l y s e s of friction stir w e l d s in single crystal a l u m i n u m and polycrystalline 2 1 9 5 a l u m i n u m reveal m a n y aspects of the textural evolution that a c c o m p a n y the microstructural development in friction stir welds. T h e s e results—extending from the initial deformation of the base plate to the subgrain d e v e l o p m e n t a n d the evolution of those subgrains into the fine, equiaxed grains near the t o o l — c a n serve a s a guide t o w a r d s understanding the textural evolution that occurs during friction stir welding of a l u m i n u m . While the results on these t w o w e l d s are consistent with e a c h other and with reports in the literature, it should also b e m e n t i o n e d that other a l u m i n u m alloys m a y be susceptible to additional m e c h a n i s m s of recrystallization. In particular, recrystallization m e c h a n i s m s that generate n e w grain orientations, such as particlestimulated recrystallization, w o u l d significantly alter these observations and m a y explain the weak texture observed in s o m e friction stir w e l d s . A schematic describing the microstructural and textural evolutions that occur ahead of the tool is s h o w n in Figure 14. T h e initial influence of the rotating F S W tool results in local rotations within the single crystal. T h e presence of small deviations from the base plate orientation c a n facilitate these rotations, leading to a divergence in the observed orientations. This orientation divergence continues to d e v e l o p with increasing temperature and deformation (closer to the F S W tool) until reaching an easily-sheared orientation with < 1 1 0 > along the shear direction and/or {111} aligned with the shear plane n o r m a l . Further rotation b e y o n d this orientation is not observed d u e to the severe decrease in rotation rate that occurs in t h e vicinity o f these orientations. T h u s , the texture evolves into a combination of the b a s e plate orientation and a terminally-rotated orientation aligned with the shear deformation. In the single crystal friction stir weld, the heat and deformation generated by t h e welding process induce static recrystallization. T h e large grains that form c o n s u m e much o f the material approaching the tool and obscure the continuous observation of texture d e v e l o p m e n t near t h e tool. T h e s e recrystallized grains b e c o m e deformed during continued w e l d i n g until they d e v e l o p
26
·
Materials Processing and Texture
T e x t u r e D e v e l o p m e n t in A l u m i n u m Friction Stir W e l d s
a texture consisting o f the B and Β c o m p o n e n t s of the simple shear texture, as is typical of hightemperature shear deformation. T h e 2195 alloy exhibits further details o f the grain and texture d e v e l o p m e n t during friction stir welding d u e to this a l l o y ' s higher resistance to recrystallization. Figure 14. Schematic of the grain d e v e l o p m e n t p r o c e s s and texture evolution occurring during friction stir welding of single crystal aluminum. R e g i o n s exhibiting each texture are indicated, as are the rotation axis, rotation directions, a n d shear plane n o r m a l direction (starred).
T h e general trend of texture development is similar in this alloy, although that development appears to arise less from the divergent rotation of slight misorientations in the b a s e plate material than from a deformationinduced subgrain formation, possibly due to a smaller initial orientation spread within the grains of the polycrystalline material. A s in the single crystal w e l d s , this w e l d exhibits subdivision of the base plate t h r o u g h subgrain formation and development. H o w e v e r , it also reveals that the process of subgrain evolution into the fine, equiaxed grains near the tool o c c u r s through a continuous d y n a m i c recrystallization m e c h a n i s m . T h e development of specific readily-sheared orientations is also observed in this weld. T h i s is evident from the relative stability of some grains while adjacent grain orientations b e c o m e subdivided into b a n d s of fine, equiaxed grains. A s in the single crystal friction stir welds, a texture consisting of t h e BIB shear c o m p o n e n t s d e v e l o p s near the tool. This BIB texture strengthens as the material is swept around the tool. Within the deposited weld, a periodic texture alternating b e t w e e n a p u r e BIB c o m p o n e n t and a mixture o f the BIB and C components is observed with a periodicity of the tool a d v a n c e per revolution. ACKNOWLEDGEMENTS T h e authors w o u l d like to dedicate this p a p e r to t h e m e m o r y of J o h n A . Wert, a great researcher, teacher, and friend. T h e authors gratefully a c k n o w l e d g e assistance from Palle Nielsen, Guilin W u , and P r e b e n O l e s e n for their assistance with the single crystal a l u m i n u m specimen. R W F a c k n o w l e d g e s support from the N a v a l Research Laboratory under the auspices of the Office of Naval Research. Support for J F B w a s provided through the Joint D o D / D O E Munitions T e c h n o l o g y D e v e l o p m e n t P r o g r a m at L o s A l a m o s N a t i o n a l Laboratory, operated by Los A l a m o s National Security, L L C under contract D E - A C 5 2 - 0 6 N A 2 5 3 9 6 . A P R acknowledges support from the National Science Foundation I / U C R C P r o g r a m Center for Friction Stir Processing ( A w a r d # E E C - 0 4 3 7 3 4 1 ) . Support for J A W w a s provided by the Danish National
Materials Processing and Texture
·
27
T e x t u r e D e v e l o p m e n t in A l u m i n u m Friction Stir W e l d s
Research Foundation t h r o u g h its support o f the Center for F u n d a m e n t a l Research: Structures in F o u r D i m e n s i o n s .
Metal
REFERENCES W . M . T h o m a s , E . D . N i c h o l a s , J . C . N e e d h a m , M . G . M u r c h , P . T e m p l e s m i t h , a n d C.J. D a w e s , "Friction Stir Butt W e l d i n g " , Int. Patent A p p . P C T / G B 9 2 / 0 2 2 0 3 a n d G B Patent A p p . 9 1 2 5 9 7 8 . 8 , Dec. 1 9 9 1 . U . S . Patent N o . 5,460,317, Oct. 1 9 9 5 . 1
H . Jin, S. Saimoto, M . Ball, a n d P.L. Threadgill, Mat Sei Tech, 1 7 , 1605-1614 (2001). Y . S . Sato, H. K o k a w a , K. Ikeda, M . E n o m o t o , S. J o g a n , a n d T. H a s h i m o t o , Met Mat Trans A, 3 2 , 941-948 (2001). D.P. Field, T . W . N e l s o n , Y. H o v a n s k i , a n d K.V. Jata, Met Mat Trans, 3 2 A , 2 8 6 9 - 2 8 7 7 (2001). K.V. Jata a n d S.L. Semiatin, Scripta Mat, 4 3 , 7 4 3 - 7 4 9 ( 2 0 0 0 ) . A . F . N o r m a n , I. B r o u g h , a n d P . B . Prangnell, Mat Sei Forum, 3 3 1 - 3 3 7 , 1713-1718 (2000). D.P. Field a n d T.W. Nelson, Mat Sei Forum, 4 0 8 - 4 1 2 , 1507-1512 (2002). Y. H o v a n s k i , T . W . N e l s o n , a n d D . P . Field, Proc. Joining o f A d v a n c e d a n d Specialty Metals. A S M Intl., Materials Park, O H , 167-171 (2001). J.-Q. Su, T.W. N e l s o n , R. M i s h r a , a n d M . M a h o n e y , Acta Mat, 5 1 , 7 1 3 - 7 2 9 (2003). M . Karlsen, 0 . Frigaard, J. Hjelen, 0 . G r o n g , a n d H. N o r u m , Mat Sei Forum 4 2 6 - 4 3 2 , 2 8 6 1 2 8 6 6 (2003). Y. Li, L.E. Murr, a n d J.C. M c C l u r e , Mat Sei Eng A, 2 7 1 , 2 2 1 3 - 2 2 3 ( 1 9 9 9 ) . J . C . M c C l u r e , E. C o r o n a d o , S. Aloor, B . M . N o w a k , L.E. Murr, and A . C . N u n e s , Jr., P r o c . 6 t h Int. Conf. o n T r e n d s in W e l d i n g R e s e a r c h . A S M Int, Materials Park, O H , p p . 2 5 7 - 2 6 1 (2003). R.W. F o n d a , J.F. Bingert, a n d K.J. Colligan, Proc. 5th Int. S y m p . o n Friction Stir W e l d i n g . T W I , C a m b r i d g e (2004). R . W . F o n d a , J.F. Bingert, a n d K.J. Colligan, Scripta Mat, 5 1 , 2 4 3 - 2 4 8 ( 2 0 0 4 ) . J.A. Schneider a n d A . C . N u n e s , Jr, Met Mat Trans Β, 3 5 B , 7 7 7 - 7 8 3 (2004). P . B . Prangnell a n d C P . H e a s o n , Acta Mat, 5 3 , 3 1 7 9 - 3 1 9 2 ( 2 0 0 5 ) . Y . S . Sato, T . W . N e l s o n , a n d C.J. Sterling, Acta Mat, 5 3 , 6 3 7 - 6 4 5 (2005). R. W . F o n d a , J. A . Wert, A . P . R e y n o l d s , a n d W . T a n g , Sei Tech Weld Join, 1 2 (4), 3 0 4 - 3 1 0 (2007). L.S. Toth, P . Gilormini, a n d J.J. J o n a s , Acta Metall, 1988. 3 6 ( 1 2 ) : p . 3 0 7 7 - 3 0 9 1 . L.S. Toth, K . W . N e a l e , a n d J.J. J o n a s , Acta Metall, 1989. 3 7 ( 8 ) : p . 2 1 9 7 - 2 2 1 0 . M.R. Barnett a n d F. Montheillet, Acta Mater, 2 0 0 2 . 5 0 ( 9 ) : p . 2 2 8 5 - 2 2 9 6 . F.J. H u m p h r e y s a n d M . Hatherly, Recrystallization a n d Related A n n e a l i n g P h e n o m e n a . P e r g a m o n , Oxford, p p . 1 7 3 a n d 3 8 8 (1995). G.R. C a n o v a , U . F . K o c k s , a n d J.J. Jonas, Acta Metall, 3 2 , 2 1 1 - 2 2 6 (1984). F. Montheillet, M . C o h e n , a n d J.J. Jonas, Acta Metall, 3 2 , 2 0 7 7 - 2 0 8 9 (1984). T. Pettersen a n d E . N e s , Met Mat Trans A, 3 4 A , 2 7 2 7 - 2 7 3 6 (2003). A . Gholina, P . Bate, a n d P. B . Prangnell, Acta Mater, 2 0 0 2 . 5 0 : p . 2 1 2 1 - 2 1 3 6 . S.C. V o g e l , D.J. Alexander, I.J. Beyerlein, M . A . M . B o u r k e , D . W . B r o w n , B . Clausen, C.N. T o m é , R . B . V o n Dreele, C . X u , a n d T . G . L a n g d o n , Mat Sei Forum, 2 0 0 3 . 4 2 6 - 4 3 2 : p . 2 6 6 1 2666. 2
3
4
5
6
7
8
9
1 0
11
1 2
13
1 4
1 5
1 6
1 7
1 8
1 9
2 0
2 1
2 2
2 3
2 4
2 5
2 6
2 7
2 8
2 9
3 0
3 1
28
R . W . F o n d a , K . E . K n i p l i n g , a n d J.F. Bingert, Scripta Mater, in press. J.H. C h o and P.R. D a w s o n , Met Mat Trans A, 2 0 0 6 . 3 7 A ( 4 ) : p . 1147-1164. T . W . N e l s o n , Personal Communication. 2007. R . W . F o n d a a n d J.F. Bingert, Scripta Mater, 5 1 , 1052-1055 (2007).
Materials Processing and Texture
MICROSTRUCTURAL STUDIES OF A FRICTION-STIR-WELDED AA5052 2
2
Ν . T. K u m b h a r ' , Santosh. K. S a h o o , R. T e w a r i ' , G. K. D e y ' , I. Samajdar , K. B h a n u m u r t h y ' Materials Science Division, B h a b h a A t o m i c research Centre, M u m b a i , India. Metallurgical Engineering and Materials Science, IIT B o m b a y , M u m b a i , India.
1
2
ABSTRACT Partially recrystallized A A 5 0 5 2 plates w e r e successfully j o i n e d through friction stir welding. C o m b i n a t i o n s of optical microscopy, microhardness, electron probe microanalysis ( Ε Ρ Μ Α ) , electron backscattered diffraction ( E B S D ) and transmission electron microscopy ( T E M ) w e r e used t o bring out a composite picture of microstructural d e v e l o p m e n t s . Different regions of the w e l d w e r e generalized as the advancing side ( A S ) , retreating side ( R S ) and nugget - using standard terminologies. A S s h o w e d evidence of shear and of grain fragmentation, while R S indicated initial stages of geometrical d y n a m i c recrystallization. The nugget region, o n the other hand, h a d nearly equiaxed grains, with strong in - grain misorientation and presence of grain interior dislocation structure ruling out contributions from static recrystallization. The nugget also had typical onion rings; each ring had a single/dominant orientation. The processes of plastic deformation, solid state-stirring and d y n a m i c recrystallization are the possible m e c h a n i s m s behind the n u g g e t microstructure. INTRODUCTION In a conventional welding p r o c e s s , a l u m i n u m alloys are typically classified as 'difficult to w e l d ' . ' T h e difficulty is often attributed to the solidification process and structure - including loss of alloying elements and presence of segregation and porosities. Friction stir welding ( F S W ) offers an alternative - an alternative through solid-state bonding, w h i c h eliminates all these problems of solidification associated with the conventional fusion welding processes. Different aspects o f microstructural d e v e l o p m e n t s ' and linking such d e v e l o p m e n t s with properties have often been the focus of t o d a y ' s research o n F S W . T h e present study m a k e s an attempt to add to the existing k n o w l e d g e - b a s e . T h e a i m of the study is to provide a composite picture of microstructural d e v e l o p m e n t associated with F S W of A A 5 0 5 2 . 11
[2
[ 3 )
Partially recrystallized A A 5 0 5 2 w a s subjected to F S W . T h e process was ' o p t i m i z e d ' to qualify radiography and die-penetrant tests - i.e. to p r o d u c e F S W w e l d s without visible m a c r o s c o p i c defects. The w e l d e d structure w a s then subjected to a variety of microstructural characterization (ranging from conventional to transmission electron microscopy, m i c r o hardness, electron p r o b e m i c r o analysis and electron backscattered diffraction), so that a composite picture of the microstructural d e v e l o p m e n t e m e r g e s . EXPERIMENTAL DETAILS 3 m m thick plates of A A 5 0 5 2 , partially recrystallized alloy with the chemical composition 2.3 M g , 0.15 Si, 0.2 F e , 0.02 Cu, 0.02 M n , 0.1 Cr, 0.01 Zn, 0.01 Ti, A l (Bai) (all in wt % ) , w e r e j o i n e d by friction stir w e l d i n g ( F S W ) . Several F S W trials were carried out. The optimised, on the basis of relative qualification through X-ray radiography and also by die penetration tests, F S W w a s performed at respective tool rotation speed of 1400 rpm, welding speed of 60 m m / m i n and tool tilt angle of 2° in the forward direction. C o m m e r c i a l high-speed steel w a s used as the tool material with shoulder diameter of 6 m m , pin diameter of 2 m m and pin height of 2.8 m m .
29
M i c r o s t r u c t u r a l S t u d i e s of a F r i c t i o n - S t i r - W e l d e d A A 5 0 5 2
T h e transverse section of the w e l d w a s used for microstructural characterization. T h e metallographically polished samples w e r e chemically etched (water: hydrofluoric acid = 100:1), so that a r o u g h estimation o f the nugget and the a d v a n c i n g and retreating s i d e s ' ' ' b e c o m e s apparent. T h e structure w a s recorded w i t h conventional m i c r o s c o p y (optical and scanning electron microscope). A c r o s s this structure, microhardness profile w a s obtained. F o r analytical microscopy, a C a m e c a S X 100 electron probe m i c r o analyzer ( Ε Ρ Μ Α ) , equipped with three wavelength dispersive spectrometers, w a s u s e d at 20 k V operating voltage. F o r electron backscattered diffraction ( E B S D ) and for transmission electron m i c r o s c o p y ( T E M ) , samples were electropolished at -40° C and at 11 volts dc in a solution containing 8 0 % ethyl alcohol and 2 0 % perchloric acid. E B S D m e a s u r e m e n t s w e r e obtained o n a Fei q u a n t a - 2 0 0 H V S E M with T S L - O I M p a c k a g e , while Jeol 3010 m i c r o s c o p e (300 k e V operating voltage) w a s used for T E M studies. 4
3
RESULTS AND DISCUSSION Figure 1(a) s h o w s the optical microstructure o f t h e F S W , while Figures 1(b) and (c) outline the respective profiles of m i c r o h a r d n e s s and Ε Ρ Μ Α line s c a n s taken across the weld. T h e regions in the F S W are classified as nugget ( N ) , advancing side interface ( A S ) , retreating side
Fig. 1—(a) Optical micrograph of the F S W specimen used in the present study. T h e regions corresponding t o nugget ( N ) , a d v a n c i n g ( A S ) and retreating ( R S ) sides are m a r k e d . A B represents the location of m e a s u r e m e n t of (b) m i c r o h a r d n e s s and (c) Ε Ρ Μ Α profiles for Ο (K„), M g (Κα) and Al (K«).
30
·
Materials Processing and Texture
M i c r o s t r u c t u r a l S t u d i e s of a F r i c t i o n - S t i r - W e l d e d A A 5 0 5 2
1
interface ( R S ) a n d base metal ( B ) - based on standard n o m e n c l a t u r e s . Relatively sharper interface between base metal a n d n u g g e t is observed on the a d v a n c i n g side than on the retreating side. A s in Figure 1(b), the n u g g e t had lower hardness t h a n the base metal - 6 3 - 7 0 and 7 0 - 7 8 V H N respectively. It is also important to note that b a s e metal across b o t h R S a n d A S h a d noticeably lower ( 5 9 ~ 6 5 V H N ) hardness. C o m p o s i t i o n profiles for Ο Κα, M g K,, a n d Al Κ , obtained from Ε Ρ Μ Α are s h o w n in Figure 1(c). T h e profiles s h o w a b s e n c e of significant oxygen p i c k - u p in the F S W , especially in the n u g g e t region. T h e profiles also indicate relatively insignificant compositional differences between different regions of the F S W . α
Fig. 2 — C o m p o s i t e image of the F S W . The figure includes an optical image ( s a m e as in Fig 1(a) outlining different regions of the F S W . C o r r e s p o n d i n g E B S D i m a g e s are also included. Figure 2 presents a c o m p o s i t e i m a g e of t h e F S W . This includes the earlier (as i n Figure 1(a)) optical image, with different regions marked. E B S D images corresponding to these regions are also included in Figure 2 . Macroscopically, or at lower magnification, the advancing side (AS) interface appeared sharper. I n an E B S D i m a g e o f higher magnification, as in Figure 3(a), the reason for this ' s h a r p n e s s ' b e c o m e s apparent. M a r k e d differences in the morphological orientation and in relative fragmentation o f the grains w e r e visible o n both sides of the interface. Grains in t h e b a s e m e t a l w e r e , in general, e q u i a x e d a n d morphologically oriented along the horizontal direction, while grains o n the other side o f the interface had 'differing' morphological orientations and also had clear p r e s e n c e of finer/fragmented grains. T h e s e differences gave the appearance of a visibly sharp interface. T h e retreating side, o n the other hand, did not present such a clear ' s w i t c h ' , see F i g u r e 3(b). Interface regions close to the retreating side, h o w e v e r , did show possibility of ' g e o m e t r i c a l d y n a m i c r e c r y s t a l l i z a t i o n ' . ' ' A s s h o w n in Figure 3(b), elongated grains o n the retreating side often had significant grain b o u n d a r y serrations - normally thought to be caused by extensive d y n a m i c r e c o v e r y . ' ' Such serrated grains are reported to get ' p i n c h e d ' , forming the basis of so-called 'geometrical d y n a m i c recrystallization'.' ' Figure 4 ( b ) strongly hints at t h e initial stages of such an event. T h e n u g g e t region, see Figure 4 ( a ) , displayed 6
7
3
Materials Processing and Texture
·
31
M i c r o s t r u c t u r a l S t u d i e s of a F r i c t i o n - S t i r - W e l d e d A A 5 0 5 2
Fig. 3 — E B S D images from (a) A d v a n c i n g side ( A S ) and (b) Retreating side (RS). In (a) morphological orientations of the grains, on the nugget side, are marked with arrows, while in ( b ) grain boundary serrations are visible. clear O n i o n r i n g s * . E a c h ring had one family of orientation (Figure 4(b)), while m a n y rings formed t h e typical * F S W shear texture (Figure 4(a)). E a c h ring of t h e ' o n i o n r i n g s ' had a single/dominant orientation, see Figure 4 ( b ) . This c a n n o t b e explained through plastic deformation alone. Plastic deformation (at least conventional plastic deformation) does not converge a relatively large n u m b e r of grains (i.e. grains/orientations corresponding to an individual ring - F i g u r e 4(a)) into a single class of orientation - and not a range of orientations or a so-called deformation fiber. A c o m b i n a t i o n of d y n a m i c recrystallization and solid-state stirring (plus c o m p l e x shear) h a d possibly contributed t o t h i s interesting microstructural feature. It is, however, fair t o admit that exact m e c h a n i s m ( s ) behind the formation of each onion ring remains largely an e n i g m a . 1
1
Fig. 4 — ( a ) E B S D i m a g e o f the nugget ( N ) . T h e image s h o w s clear patterns of ' o n i o n r i n g s ' also m a r k e d by d r a w n lines. E a c h ring had o n e type of orientation, while the entire n u g g e t h a d a typical F S W shear texture. T h e s e are s h o w n in the pole figures (b) for entire region and (c) for o n e 'onion ring'. P W D and T D in the pole figures respectively represent perpendicular to welding direction and transverse direction.
32
·
Materials Processing and Texture
M i c r o s t r u c t u r a l S t u d i e s of a F r i c t i o n - S t i r - W e l d e d A A 5 0 5 2
Table I: Table s u m m a r i z i n g t h e E B S D estimated microstructural features corresponding to the different regions of the F S W . T h e regions (B - b a s e metal, A S - a d v a n c i n g side, Ν - nugget and R S - retreating side) are marked in Figure 3 . Regions
A v e r a g e grain diameter (microns)
Texture index
Grain average orientation spread (GOS) (degrees)
Grain average aspect ratio
Β
15.95
102
1.27
0.48
AS
5.55
125
1.48
0.45
Ν
5.09
142
1.05
0.50
RS
9.80
95
2.24
0.47
Table I s u m m a r i z e s the microstructural observations at different regions o n the F S W . T h e s u m m a r y (Table I) or t h e details include a r a n g e of parameters - grain size and grain aspect ratio, grain orientation spread ( G O S ) a n d texture index values. It is important to note that the nugget region had clear refinement in grain size and also had presence of nearly equiaxed grains (aspect ratio ~ 0.5). T h e G O S values w e r e also lower, in general, in the nugget. The combination indicates possible involvement of both plastic deformation and softening in the formation of the nugget. It n e e d s to be pointed out that the lowest misorientation values w e r e recorded (see Table 1) o n the base metal side of the R S and A S interfaces - an observation w h i c h is in agreement with the patterns of microhardness profile (see Figure 1(b)). H e r e it is important to and thermal discuss/hypothesize about the t w o apparent causes - plastic d e f o r m a t i o n ' activation t h r o u g h localized heating/friction.' ' to explain the observed patterns of microstructural d e v e l o p m e n t s . Indirect evidence o f thermal activation is identifiable from clear softening o f the partially recrystallized base m e t a l near the F S W interfaces, see Figure 1(b). ,0]
11
Fig. 5—Bright field T E M i m a g e of the n u g g e t region showing grains with dislocation substructure inside.
Materials Processing and Texture
•
33
M i c r o s t r u c t u r a l S t u d i e s of a F r i c t i o n - S t i r - W e l d e d A A 5 0 5 2
Direct evidences of thermal activation vs. plastic deformation are difficult t o distinguish or demarcate - their effects o n the respective microstructural parameters (grain interior dislocation structure, grain refinement, etc.) often b e i n g cumulative and/or subtractive. T o explore the microstructural d e v e l o p m e n t s in the nugget further, T E M samples w e r e studied. A s s h o w n in F i g u r e 5, the equiaxed nugget grains had clear signatures of plastic deformation - signatures both in t e r m s of dislocation substructure and also in misorientation build-up. Microstructural evidence, especially the T E M evidence (see Figure 5) o n in-grain dislocation substructure and o n misorientation build-up, rule out possible role of static recrystallization. Roles of recovery ( d y n a m i c and possibly static) and d y n a m i c recrystallization are, however, apparent - latter accounting for the equiaxed grain shape of the nugget z o n e . CONCLUSIONS F S W trials w e r e conducted o n partially recrystallized A A 5 0 5 2 . Analytical microscopy established chemical h o m o g e n e i t y and absence of o x y g e n pick-up. A S s h o w e d evidence of shear and of grain fragmentation, while RS indicated initial stages of geometrical d y n a m i c recrystallization. T h e n u g g e t region had nearly e q u i a x e d grains, with strong in-grain misorientation and presence of grain-interior dislocation structure ruling out contributions from static recrystallization. T h e nugget also had typical onion rings with each ring having a single/dominant orientation. T h e process of plastic deformation, solid-state stirring and d y n a m i c recrystallization are the suspected m e c h a n i s m s behind the nugget microstructure. T h e F S W microstructure d e v e l o p m e n t w a s clearly asymmetric. In terms of visible microstructure and/or different microstructural indices (e.g. grain size refinement, misorientation d e v e l o p m e n t , etc.), strongest microstructural d e v e l o p m e n t was in the nugget region followed by A S and t h e n by R S . REFERENCES C.J. D a w e s and W . M . T h o m a s , Weld. J., 7 5 , 41-45 (1996) Y . S . Sato, H. K o k a w a , M . E n o m o t o and S. Jogan, Metall. Mater. Trans. Α., 3 0 , 2 4 2 9 - 2 4 3 7 (1999) W . M . T h o m a s and E . D . N i c h o l a s , Mater. Des., 1 8 , 269-273 (1997) M . Guerra, C. Schmidt, J.C. M c C l u r e , L.E. Murr and A . C . N u n e s , Mater Char., 4 9 , 95-101 (2003) R . W . F o n d a , J.F. Bingert and K. J. Colligan, Scripta Mater., 5 1 , 2 4 3 - 2 4 8 (2004) H. Z h a n g , S.B. Lin, L. W u , J.C. Feng, Sh.L. M a , Mater. Des., 2 7 , 8 0 5 - 8 0 9 (2006) M . E . K a s s n e r and S.R.Barrabes, Mater. Sei. Eng. A, 4 1 0 - 4 1 1 , 1 5 2 - 1 5 5 (2005) K . N . Krishnan, Mater. Sei. Eng. A, 3 2 7 , 2 4 6 - 2 5 1 (2002) R . D . Doherty, D.A. H u g h e s , F.J. H u m p h r e y s , J.J. J o n a s , D. Juul Jensen, M.E. Kassner, W.E. King, T.R. M c N e l l e y , H.J. M c Q u e e n and A . D . Rollett, Mater. Sei. Eng. A, 2 3 8 , 2 1 9 - 2 7 4 (1997) T. W a t a n a b e , H. T a k a y a m a and A . Yanagisawa, J Mater. Process. Technol, 1 7 8 , 3 4 2 - 3 4 9 (2006) L.E. Murr, A. Trillo, S. Pappu, C. K e n n e d y , J. Mater. Sei, 3 7 , 3 3 3 7 - 3 3 6 0 (2002) 1
2
3
4
5
6
7
8
9
1 0
11
34
Materials Processing and Texture
EFFECT OF MICROSTRUCTURE ON MECHANICAL PROPERTIES OF FRICTIONW E L D E D J O I N T S B E T W E E N A L U M I N U M A L L O Y S ( 6 0 6 1 , 5052) A N D 3 0 4 S T A I N L E S S STEEL 1
1
1
1
2
2
Ν . T. K u m b h a r , A. L a i k , G. K. D e y , Κ. B h a n u m u r t h y , J. K r i s h n a n , D. J. D e r o s e , R. L. Suthar , S. K. S a h o o and I. S a m a j d a r Materials Science Division, B h a b h a A t o m i c R e s e a r c h Centre, M u m b a i , I n d i a - 4 0 0 0 8 5 . Centre for Design and Manufacture, B h a b h a A t o m i c Research Centre, M u m b a i , I n d i a - 4 0 0 0 8 5 . Metallurgical Engineering and Materials Science, IIT B o m b a y , M u m b a i , I n d i a - 4 0 0 076. 2
3
3
1
2
3
ABSTRACT C o m m e r c i a l A l u m i n u m alloys 5 0 5 2 and 6061 w e r e welded to A I S I 3 0 4 stainless steel using a continuous drive friction welding technique. In order to study the effect of interlayer on the j o i n i n g efficiency, C u interlayer w a s diffusion b o n d e d to stainless steel and subsequently friction welded to Al alloys. T h e microstructure, microtexture and mechanical properties were investigated and process parameters w e r e optimized. T h e reaction layer essentially consisted of FeAl intermetallics and this layer formed with 2 μπι thickness under the condition of relatively long friction time ( t l ) and l o w upset pressure (P2). T h e interface b e t w e e n A l alloy/ stainless steel 304 w a s very sharp for direct friction welded specimens and the interfaces at stainless steel 304/Cu and Cu/Al alloy for specimens welded with C u interlayer were diffuse in nature. Orientation imaging m i c r o s c o p y results for direct friction w e l d e d s p e c i m e n showed that the average grain size of Al close to t h e interface w a s around 100 urn and the extent of misorientations at the interface w a s found to be m a x i m u m . T h e p a p e r presents the details of the microstructure evolved at the interface during friction welding and also attempts to correlate with the observed mechanical properties. Key words: Friction welding, diffusion welding, 6061 Al alloy, 5 0 5 2 Al alloy, and Mechanical properties. INTRODUCTION T h e r e is an increasing need for j o i n i n g Al alloys to stainless steel for several areas including cryogenic applications, h i g h v a c u u m c h a m b e r s and detector systems. Diffusion bonding technique h a s been extensively used for j o i n i n g planar a s s e m b l y and this technique has yielded excellent mechanical properties for j o i n i n g Al/stainless steel specimens " l Friction welding is an inevitable technique for j o i n i n g tubular assemblies and this technique is used for j o i n i n g stainless steel , Ti alloys ' and also Al alloys ' . The p r o b l e m s associated with the welding of dissimilar materials are essentially due to difference in hardness, melting points and also due to the occurrence of solid state reactions locally at the interface. The success of the joints by friction welding strongly d e p e n d s on the process parameters. T h e formation of brittle intermetallics at the interface for Al/stainless steel j o i n t d e p e n d s mostly o n the friction time, the duration of t i m e for relative m o v e m e n t of parts under frictional forces. H o w e v e r , for short time under low friction pressure and low upset pressures, the j o i n t has poor strength. In order to achieve g o o d m e c h a n i c a l properties, it is necessary to w o r k at short friction time and high friction and upset pressure with out sacrificing t h e A l parent material. In order to m i n i m i z e the flashing of Al and also to optimize the process parameters in a broader w i n d o w , it is important to use metallurgically compatible interlayer with stainless steel and a l u m i n u m . Several studies have been reported o n the friction welding of a l u m i n u m to stainless steel l Initial w o r k on friction joining w a s mostly on the process parameter and evaluation of mechanical properties " ' . [1
[3)
[4
2
|5
| 6
[7
13
35
Effect of M i c r o s t r u c t u r e o n M e c h a n i c a l P r o p e r t i e s of F r i c t i o n - W e l d e d J o i n t s
H o w e v e r the correlation of the evolved microstructure, mechanical properties with process parameters is not well studied for 6061 Al alloys to stainless steel 304 j o i n t s . EXPERIMENTAL WORK C o m m e r c i a l l y available 5 0 5 2 Al alloy, 6061 Al alloy and 304 austenitic stainless steel in the form of r o d s (30 m m dia and 4 0 m m long) w e r e used in these studies. T h e chemical composition of these alloys are Al 6061 [0.4-Si, 0.7-Fe, 0.15-Cu, 0.15-Mn, 0.8-Mg, 0.04-Cr, 0.5Zn, 0.15-Ti, A l - B a l ] , Al 5 0 5 2 [0.25-Si, 0.4-Fe, 0.1-Cu, 0.1-Mn, 2.2 M g , 0.15-Cr, 0.1-Zn, AlBal], stainless steel 304 [0.07-C, 0.3-Si, 1.38-Mn, 0.4-P, 0.28-S, 18.4-Cr, 8.1 N i , F e - B a l ] . Both Al alloys used w e r e received in solutionised condition and this condition w a s specially chosen for optimizing the process parameter especially friction pressure Pi and upset pressure P'2 at lower values. Friction w e l d i n g of Al alloy to steel w a s carried out in t w o different w a y s . Al alloy w a s directly b o n d e d to stainless steel and the details of variation in the process parameters are listed in Table II. T h e second m e t h o d involoves, diffusion b o n d i n g of steel to copper with N i as an interlayer (10-15 p m ) and subsequently the friction welding w a s carried out b e t w e e n the diffusion bonded a s s e m b l y with Al alloy. T h e details of the process parameters are listed in Table II. For all experiments the r p m w a s kept at 1980 and upset t i m e t2 w a s fixed at 5 sec. T h e cross sections of the j o i n t s w e r e polished with 1 μηι d i a m o n d paste and the interface w a s electrochemically etched in oxalic acid. T h e microstructure of interface w a s e x a m i n e d by optical microscopy and also by S E M - E D X . For electron backscattered diffraction ( E B S D ) , samples w e r e electropolished at -40° C and at 11 volts dc in a solution containing 8 0 % ethyl alcohol and 2 0 % perchloric acid. E B S D m e a s u r e m e n t s w e r e obtained on a Fei q u a n t a - 2 0 0 H V S E M with T S L - O I M p a c k a g e . T h e joint strength w a s evaluated by tensile testing (gauge length 3 0 m m and g a u g e dia 6 m m ) at r o o m temperature using Instron w i t h a crosshead speed of 0 . 5 m m /min. In addition the quality of the joint w a s also checked by helium leak testing. Table I - Mechanical property of parent materials Yield strength ( M P a )
UTS (MPa)
Hardness (VHN) 80
A15052
90
195
A16061
276
230
98
SS304
205
515
153
Cu
69
220
50
Table II - Material c o m b i n a t i o n s and the process parameters used for friction welding. All the trials w e r e d o n e at 1980 r p m and upset time t 5 sec. 2
36
Material combination
Inter layer
Friction Pressure Pi (MPa)
Upset Pressure P (MPa)
2
Friction T i m e ti (Sec) 0.5-2
1
SS 3 0 4 - A l 6061
Direct
40-70
100-140
2
S S 3 0 4 - A l 6061
Cu
40-70
100-140
0.5-2
3
SS 3 0 4 - A l 5052
Direct
40-70
100-140
0.5-2
4
SS 3 0 4 - Al 5 0 5 2
Cu
40-70
100-140
0.5-2
·
Materials Processing and Texture
Effect of M i c r o s t r u c t u r e o n M e c h a n i c a l P r o p e r t i e s of F r i c t i o n - W e l d e d Joints
RESULTS A N D DISCUSSION Microstructure of the interface for the direct b o n d e d a l u m i n i u m alloy / stainless steel (Al alloy/stainless steel). Secondary Electron micrographs ( S E M ) for direct friction welded stainless steel 3046061 A l alloy a r e s h o w n in F i g . l (a) and S E M taken for the j o i n t interface at higher magnification is s h o w n in F i g . l (b). Fig 1(a) s h o w s excellent contact all along the interface. D u e to heavy etching on the Stainless Steel small gap is seen and for all unetched specimens this is no debonding at the interface. This specimen w a s processed at optimized Friction pressure (Pi) = 60 M P a , Upset pressure (PÎ)=140 M P a , friction time (ti)=l sec and upset time t2 — 5 sec. (see Table HI). The formation of intermetallic clearly evident in F i g . l (b). It can be seen that the width of this intermetallic w a s less than 2 um. Width o f this intermetallic increased with increase in ti and this had adverse effect on the mechanical properties (discussed in the next section). E D X analysis of this layer w a s close t o Fe AI (Cr, N i , Ο . M g ) . Earlier studies by Sundaresan and Murthi et al ' o n commercially pure Al/stainless steel h a v e reported the formation of Fe2Als and FeAl3 and the studies by F u k u m a t o et al ' h a v e s h o w n the formation of Fe2AI , FeAl and F e A l . T h e formation of these c o m p o u n d s strongly d e p e n d s o n the local temperature attained during welding and the formation o f F e A l (O) in the present experiments closely matches with the w o r k of F u k u m a t o ' . Similar microstructural features w e r e observed for direct friction welded stainless steel 3 0 4 - 5052 Al alloy specimen. 6)
1 4 1
5
3
f3
Fig.l (a) Microstructure o f the direct friction welded stainless steel 3 0 4 / A l 6061 specimen (b) Microstructure o f the direct friction welded stainless steel 3 0 4 / A l 6061 specimen at higher magnification. Fig 2 s h o w s the inverse pole figure m a p o f the direct friction welded stainless steel/Al 5052 alloy. In the above figure it is seen that near the friction welded interface on the A A 5052 side, there is considerable a m o u n t of grain size refinement, whereas there i s n ' t any on the stainless steel 304 side. T h e A l alloy grains are refined upto a range of about 100 μπι beyond which the grains size is m o r e or less equal to that of the parent material. A very fine layer of intermetallics o f AI2CU and A L C u (O, M g ) can be observed w h i c h is the darker (nonindexed)
Materials Processing and Texture
·
37
Effect of M i c r o s t r u c t u r e o n M e c h a n i c a l P r o p e r t i e s of F r i c t i o n - W e l d e d J o i n t s
Fig. 2 Inverse pole figure m a p of direct friction welded stainless steel 304 / Al 5 0 5 2 portion between the stainless steel and a l u m i n i u m alloy. It is observed that the average grain size the stainless steel 3 0 4 is around 8 μπι and that in the Al 5 0 5 2 alloy, increases gradually from 4 urn to 20 urn within the first 80 μπι beyond w h i c h it saturates to that of the parent material. T h e grain average misorientation ( G A M ) values are t h e highest very near the friction w e l d i n g interface and decrease rapidly within a range of 50 μπι to around a value of 0.3 on the S S 3 0 4 side, w h e r e a s such a rapid fall is not o b s e r v e d o n the Al side. T h e variation of Kernel average misorientation ( K A M ) is similar t o that of the G A M for t h e SS 3 0 4 side. But o n the Al 5052 side these values are constant upto 60 μηι beyond which there is a gradual decrease to value of 0.5 from 2 0 0 μηι and b e y o n d . Microstructure o f the friction welded using copper as an interlayer (stainless steel/Cu/aluminium alloy). Secondary Electron M i c r o g r a p h s ( S E M ) for stainless steel 3 0 4 / C u and C u /Al 5 0 5 2 specimen under optimized condition (see Table III) are shown in Fig.3 (a) and Fig.3 (b). Fig.3 (a) s h o w s excellent contact all a l o n g t h e interface. A s discussed in earlier, stainless steel / C u specimen is essentially diffusion b o n d e d and the p r e s e n c e of Kirkendall m i c r o p o r e s o n the C u side is essentially d u e to faster mobility o f Cu. T h e presence of electrodeposited Ni layer of about 10 μπι is clearly evident at stainless steel/Ni interface. In addition, solid solution of N i (Cu) of about 10 μπι width is also evident o n the C u rich side. These regions are m a r k e d o n this m i c r o g r a p h . T h e interface correspond to Cu/Al is s h o w n in Fig.3 (b). The reaction zone of width 10-15 μπι can be seen and m i c r o g r a p h taken at the higher magnification s h o w e d debonding at the reaction
Fig.3 Secondary Electron M i c r o g r a p h of (a) stainless steel 304 / C u interface and (b) C u /Al 5 0 5 2 interface.
38
Materials Processing and Texture
Effect of M i c r o s t r u c t u r e o n M e c h a n i c a l P r o p e r t i e s of F r i c t i o n - W e l d e d Joints
zone/Al interface. T h e s e s p e c i m e n s w e r e friction w e l d e d at friction t i m e ( t i ) = l sec. T w o intermetallic c o m p o u n d s can be noticed in Fig.3 (b) w h i c h s h o w s an intermetallic of less than 2 μπι close t o C u layer and a reaction z o n e of about 10 μπι o n A l side. T h e composition of these intermetallics are m e a s u r e d by E D X analyses and found to be A l C u and A h C u ( O , M g ) . The nature o f microstructure for stainless steel/Cu/6061 Al w a s found to b e similar to that of stainless steel/Cu/5052 Al expect that the w i d t h of the reaction zone w a s found to be less than 5 μπι. The formation o f t h e s e intermallics is basically responsible for brittle failure o f these joints. 2
Fig 4 - Inverse pole figure (IPF) m a p s of (a) diffusion b o n d e d stainless steel 304-Cu and (b) friction w e l d e d 6061 Al. Fig 4 a b o v e s h o w s the inverse pole figure m a p s of (a) the interface between diffusion bonded stainless steel 304 and copper and (b) t h e friction welded interface b e t w e e n copper and A A 6 0 6 1 . In the above figures the region corresponding to the copper interface could not be indexed properly d u e to heavily deformed grains (with very l o w confidence index) and hence appears t o be darker. T h e nature and distribution of the grain size, grain average misorientation and kernel average misorientation o n the a l u m i n i u m side is nearly the s a m e as that of the direct friction welded s p e c i m e n s . Mechanical Properties T h e m i c r o h a r d n e s s profiles taken across the b o n d e d interface for all the four geometries at optimized condition (Table III) are s h o w n in Fig.5 (a),(b) and (c). T h e hardness value o n the Al side is in t h e r a n g e of 50 V H N and on the stainless steel side is around 2 0 0 V H N . T h e microhardness value for direct w e l d e d specimen, (Fig.5 (a)), s h o w a hardness value o f about 238 V H N at t h e interface. T h e m i c r o h a r d n e s s value for Al 6 0 6 1 / C u and A l 5052/Cu close to the interface is generally less t h a n 150 V H N (Fig.5 (b) and (c)) and for Cu/stainless steel interface this value is close to 2 5 0 V H N (Fig.5 (b) and 5(c)). It can be seen that substantial hardening is taking place at all the faces m e n t i o n e d a b o v e . The increase in the h a r d n e s s value for the stainless steel side seems to be h i g h . T h e origin of large increase in t h e m i c r o h a r d n e s s values at the interface could b e attributed to a) high plastic deformation at the interface & b ) formation of intermetallics. D e p e n d i n g o n the degree of plastic deformation occurring at the interface, it is
Materials Processing and Texture
·
39
Effect of M i c r o s t r u c t u r e o n M e c h a n i c a l P r o p e r t i e s of F r i c t i o n - W e l d e d J o i n t s
A 6061
Cu
, SS304
D 0
500
Distance (microns)
1000
1500
2000
2500
300C
1000
1500
Distance (microns)
Distance (microns)
(a)
(b)
(c)
Fig. 5 Vickers hardness Al alloy/ S S 3 0 4 and A l alloy/Cu/SS 304 near weld interface observed that there is decrease in the grain size and increased high defect density, resulting in higher hardness. Tensile test results for all the geometries under optimized condition are s h o w n in Table III. These results clearly s h o w that the direct welded stainless steel 304 / 5052 A l and stainless steel 304/6061 Al samples s h o w superior mechanical properties and mostly failed in t h e parent material. O n the other hand the specimen stainless steel 3 0 4 / C u / 5 0 5 2 Al failed at Cu/Al interface and the strength of the joint is 130 M P a close to the parent Al material. Least b o n d strength w a s obtained for stainless steel 304/Cu/6061 Al and this joint failed also at Cu/Al interface and the strength is only 100 M P a , m u c h lower t h a n the U T S of 6061 A l (230 M P a ) All the specimens passed helium test and these results confirm t h e complete attainment of contact all along the interface Table III. Mechanical Property under optimized friction welding condition Sr. No.
Process parameter
Material combinations
tl
Pi
(MPa)
(MPa)
(sec)
Mechanical properties of the weld
t2
Failure location
(sec)
Strength MPa
Elongation
135
13
5x10-'°
2
4x10-'°
13
8x10-'°
2
5x10·'°
1
SS304-A15052
35
140
1
5
Parent material (Al)
2
SS304-CU-A15052
35
140
1
5
Failed at Cu-AlInterface
130 228
3
SS304-A16061
60
140
1
5
Parent material (Al)
4
SS304-Cu-A16061
60
140
1
5
Failed at Cu-AlInterface
Helium leak testing mbar l/s
100
%
CONCLUSIONS A l u m i n i u m alloys 5 0 5 2 Al and 6061 Al alloys w e r e b o n d e d directly to stainless steel 3 0 4 and also using C u interlayer b y continuous drive friction welding technique. T h e processes o f welding and optimizing of their parameters w e r e investigated by microstructural observation, microhardness m e a s u r e m e n t and also by tensile testing. The important conclusions based o n the experimental result are s u m m a r i z e d below:
40
Materials Processing and Texture
Effect of M i c r o s t r u c t u r e o n M e c h a n i c a l P r o p e r t i e s of F r i c t i o n - W e l d e d Joints
T h e direct friction w e l d e d stainless steel 3 0 4 and A A 5052 s h o w e d that there exists an extremely deformed r e g i o n o f approximately 100 μηι w i t h very fine grains o n the aluminium side. N o such refined zone is o b s e r v e d o n the stainless steel side. A t regions very close to the friction welding interface the extent of misorientation (grain a v e r a g e misorientation and kernel a v e r a g e misorientation) is the highest both in the stainless steel side as well as the a l u m i n i u m side, while it d r o p s on either side o f the interface. T h e s a m e is true for the s p e c i m e n s w h i c h w e r e friction w e l d e d using c o p p e r as an interlayer. Direct b o n d e d s p e c i m e n s s h o w e d superior m e c h a n i c a l properties c o m p a r e d to specimen using C u interlayer. T h e j o i n t strength increased with increase in friction t i m e t l . A t shorter time sound j o i n t s could n o t obtained, and for longer time the thickness o f the intermetallic c o m p o u n d s is increased, resulting in brittle failure o f the joint. A n o p t i m u m time ti =1 sec s e e m s to be suitable for direct friction welding. REFERENCES 1
2
3
4
5
6
7
8
9
1 0
1 1
1 2
1 3
1 4
K. B h a n u m u r t h y , R.K. Fotedar, D.J. Derose, G . B . K a l e , A.L. P a p p a c h a n , A.K. Grover, J. Krishnan, Mater Sei and Tech, 2 2 , (2006) J. K r i s h n a n , K. B h a n u m u r t h y , P. S. G a w d e , D . J. D e r o s e , G. B . K a l e and G. Krishnamurthy, J Mater Pro Tech, 6 6 , 8 5 - 8 9 , ( 1 9 9 7 ) S. F u k u m o t o , H. T s u b a k i n o , O . Okita, M . Aritosh a n d T. Tomita, Mater Sei and Tech,.lS, 1080-1086,(1999) W o n - B a e L e e and S e u n g - B o o J u n g , Materials Transaction, 4 5 , 2 8 0 5 - 2 8 1 1 , (2004) T. M o h a n d a s , D . Banerjeee, V . V . K u t u m b a R a o , Mater Sei andEnggA., 2 8 9 , 7 0 - 8 2 , (2000) S. Sundaresan, K.G.K. Murti, Material Forum, 1 7 , 3 0 1 - 3 0 7 , (1993) Bekir S. Yilbas, A . Z . Sahin, Nafiz K a h r a m a n , A h m e d z Al-lGarni, J Mater Pro Tech, 4 9 , 431-443,(1995) O.Yilmaz, H . Celik, J Mater Pro Tech, 1 4 1 , 6 7 - 7 6 , ( 2 0 0 3 ) M u m i n Sahin, H Erol A k a t a , J Mater Pro Tech, 1 4 2 , 2 3 9 - 2 4 6 , (2003) S. F u k u m o t o , H . T s u b a k i n o , K.Okita, M . A Ritoshi and T. T o m i t a , Mater Sei and Tech, 1 4 , 3 3 3 - 3 3 8 , (1998) N . Ö z d e m i r , Materials Letters, 5 9 , 2 5 0 4 - 2 5 0 9 , (2005) V . V . Satyanarayana, G. M a d h u S u d h a n R e d d y , T. M o h a n d a s , J Mater Pro Tech, 1 6 0 , 1 2 8 131,(2005) M . Y i l m a g , M . Col, M . Acet, Material Characterization, 4 9 , 4 2 1 - 4 2 9 , (2003), S. F u k u m o t o , H . T s u b a k i n o , K.Okita, M . A Ritoshi and T. T o m i t a , Mater Sei and Tech, 1 3 , 6 7 9 - 6 8 6 , (1997)
Materials Processing a n d Texture
41
This Page Intentionally Left Blank
EFFECT OF T E X T U R E O N F R A C T U R E LIMIT STRAIN IN FRICTION STIR PROCESSED AZ31BMG ALLOY Y u t a k a S. Sato, Akihiro Sugimoto, Hiroyuki K o k a w a T o h o k u University Department of Materials Processing, Graduate School of Engineering, 6-6-02 Aramaki-azaAoba, Aoba-ku Sendai 9 8 0 - 8 5 7 9 , Japan Chang-Woo Lee K o r e a Institute of Industrial T e c h n o l o g y Production T e c h n o l o g y Center, 9 9 4 - 3 2 , D o n g c h u n - D o n g , Y e o n s u - G u Incheon 4 0 6 - 1 3 0 , South K o r e a
ABSTRACT M g alloy h a s p o o r ductility and formability at r o o m temperature because of lack of the active slip systems. Recently, it h a s b e e n reported that the grain refinement improves its ductility, thus it w o u l d be expected that friction stir processing (FSP) enhances the formability of M g alloys because F S P can create h o m o g e n e o u s microstructure consisting of fine grains. H o w e v e r , F S P leads to not only the grain refinement, but also the strong accumulation of crystallographic texture in M g alloy. It is likely that t h e texture formed during F S P also affects the formability significantly in M g alloy, but effect of texture on formability in FSPed M g alloy has not been e x a m i n e d yet. In this study, F S P w a s applied to A Z 3 1 B M g alloy at several FSP parameters and the stir zones having the different microstructural features w e r e obtained. Effect of microstructure, especially grain size and texture, o n fracture limit major strain under the uniaxial tension w a s e x a m i n e d in the stir zone of the F S P e d A Z 3 1 B M g alloy, and the microstructural features for the formability e n h a n c e m e n t in M g alloy w e r e proposed. INTRODUCTION M g alloys h a v e a hexagonal close-packed (h.c.p.) crystallographic structure. In M g alloys, only the basal slip system is active at r o o m temperature, because the critical resolved shear stresses of the basal slip system are about one hundredth of those of non-basal slip systems. Since the basal slip can provide only t w o independent slip systems which are insufficient for plastic deformation of the polycrystal, M g alloys exhibit poor ductility and formability at r o o m temperature. Recently, s o m e p a p e r s " h a v e reported that high ductility is obtained in fine grained M g alloys. Koike et a l . clarified that the i m p r o v e m e n t of ductility by t h e grain refinement was due to activity of non-basal slip systems and grain b o u n d a r y sliding induced by plastic compatibility stress associated with grain boundaries, a n d d y n a m i c recovery. T h e s e studies suggest that the grain refinement can raise formability in M g alloys. Friction stir processing ( F S P ) , a d e v e l o p m e n t based on friction stir welding ( F S W ) , is a solid-state processing technique used for the microstructural modification. This process is often utilized to refine the grain structure or to h o m o g e n i z e the heterogeneous microstructure . It is well k n o w n that F S P can create fine recrystallized grains in several kinds of metallic materials. " 1
2
4
4
5
5
6
7
43
Effect of T e x t u r e o n F r a c t u r e Limit S t r a i n in Friction Stir P r o c e s s e d A Z 3 1 Β M g Alloy
Grain size of the m o s t recrystallized grains lies b e t w e e n 1 and 20 μηι in the stir zone of Al and M g a l l o y s , " " w h i c h w o u l d be small e n o u g h to enhance the formability in M g alloys. According to this background, the present authors applied multi-pass F S P to A Z 9 1 D and A Z 3 1 B M g alloys to enhance formability of M g alloys, and reported in previous p a p e r s ' ' that F S P significantly raised the formability in M g alloys, and that the finer grains resulted in the higher formability. H o w e v e r , F S P leads to not only the grain refinement, but also the strong accumulation of crystallographic texture, associated with material flow, in M g a l l o y . ' Since the M g alloys exhibit a strong plastic a n i s o t r o p y , " the texture formed during F S P should also affect the formability significantly in M g alloys, but effect of texture on formability in F S P e d M g alloy h a s not been e x a m i n e d yet. In the present study, the stir zones having the different microstructural features were produced in A Z 3 1 B M g alloy by F S P at several F S P parameters. T h e microstructural features and fracture limit major strain under the uniaxial tension w e r e examined. T h e objective of this study is to clarify effect of microstructure, especially grain size and texture, o n ductility in the FSPed M g alloy systematically. 7
2
3
4
EXPERIMENTAL PROCEDURES T h e base material used in this study w a s a hot-rolled A Z 3 1 B M g alloy ( 3 w t % Al and l w t % Z n ) , 2 m m in thickness. F S P w a s applied to this alloy using a general F S P tool. Shoulder diameter, pin diameter and pin length of the F S P tool were 15 m m , 6 m m and 1.7 m m , respectively. F S P w a s performed at several F S P parameters. Variables used are the rotational speed, the travel speed, the plunge depth and the tilt angle (tool-to-workpiece angle from the vertical axis) of the F S P tool. C o m b i n a t i o n s of the parameters e m p l o y e d in this study are s h o w n in Table I. Table I. FSP Parameters U s e d in This Study Parameter N o . P l u n g e depth Title angle (mm) (deg) 1 1.8 3 1.7 3 2 3 1.9 3 4 1.8 1 1.8 5 5 1.8 6 3 7 1.8 3 1.8 8 3 9 1.8 3 10 1.8 3 11 1.8 3 12 1.8 3 1.8 3 13
Rotational speed (rpm) 500 500 500 500 500 400 750 1000 1500 2000 2600 500 500
Travel speed (mm/s) 12 12 12 12 12 12 12 12 12 12 12 5.7 1.45
After F S P , sample for optical microscopy w a s cut perpendicular to the F S P direction from the FSPed alloy. The sample surface w a s electrolytically polished in a 6 ml perchloric acid + 25 ml butoxyethanol + 4 0 0 ml ethanol solution at 223 K, and then etched in a 5 m l acetic acid
44
·
Materials Processing and Texture
Effect of T e x t u r e o n F r a c t u r e Limit S t r a i n in Friction Stir P r o c e s s e d A Z 3 1 Β M g Alloy
+ 6 g picric acid + 10 m l w a t e r + 100 ml methanol solution. Grain size in the stir zone w a s measured b y m e a n linear intercept m e t h o d for optical micrographs. The crystallographic texture w a s investigated by taking X-ray diffraction pole figures. A Shultz reflection m e t h o d w a s e m p l o y e d in an a z i m u t h angle r a n g e of 0 - 85 d e g with respect to the normal direction of the F S P e d alloy. N o background correction w a s m a d e in the pole figures. Schematic illustration of s a m p l e locations for t h e texture analysis is d r a w n in Fig. 1. 5 samples, i.e. the upper region o n t h e centerline ( C U ) , the m i d d l e region o n the centerline ( C C ) , the bottom region on the centerline ( C B ) , the retreating side o n the midsection (RS), and the advancing side on the midsection ( A S ) , w e r e cut from the stir z o n e by electrical discharge m a c h i n e ( E D M ) . (0002) pole figure w a s obtained from the 5 regions of the stir z o n e , as s h o w n in Fig. 1. Coordinate a x e s for pole figures are the F S P direction ( F S P D ) , transverse direction ( T D ) and n o r m a l direction ( N D ) of the plate (Fig. 1).
F i g u r e 1. Schematic illustration of sample locations for t h e texture analysis.
Figure 2 . Configuration of uniaxial tensile specimen used in this study. Fracture limit major strain under t h e uniaxial tension w a s evaluated. Configuration of the test s p e c i m e n s are illustrated in Fig. 2. Specimen for the uniaxial tension w a s cut from only the stir zone by E D M , and t h e n inequalities of the upper and lower surfaces were eliminated in order to only e x a m i n e t h e effect of microstructure o n the strain. After t h e elimination, paint w a s sprayed to the specimen surface, and m a n y spots o f the paint w e r e put on the surface, as shown
Materials Processing and Texture
·
45
Effect of T e x t u r e o n F r a c t u r e Limit Strain in Friction Stir P r o c e s s e d A Z 3 1 Β M g Alloy
in Fig. 3. After taking a picture of t h e surface with m a n y paint spots, the uniaxial tensile test w a s conducted. T h e fracture limit major strain (5m#or) w a s m e a s u r e d using the following equation a n d distance b e t w e e n t w o spots closest to the fracture location after tensile test. Emajor = Ή ( / / l(, )
(1)
where and / are distances b e t w e e n t w o spots along the tensile direction before and after test, respectively. Tensile test w a s carried out at room temperature at a cross-head speed of 1.67 χ 10" mm/s. 2
Figure 3 . M a c r o s c o p i c o v e r v i e w of the tensile specimen with sprayed paint spots.
Figure 4 . (a) Cross section o f the FSPed alloy produced using P a r a m e t e r 1, and Optical micrographs of the b a s e material ( B M ) region (b), transition z o n e ( T Z ) (c), a n d stir zone (SZ) (d) of this FSPed alloy. RESULTS AND DISCUSSION Cross section perpendicular to the F S P direction of the alloy is shown in Fig. 4(a). T h i s F S P e d alloy was p r o d u c e d using Parameter 1. T h e stir zone d o e s not have any porosity and defects. Optical m i c r o g r a p h s o f the stir zone (SZ), the transition zone ( T Z ) located outside t h e S Z , and the unaffected b a s e material ( B M ) region are given in Fig. 4(b)-(d), respectively. T h e B M has the coarse grain structure including twins. A v e r a g e grain size of the B M w a s 18 μ ι η . T h e SZ has the fine equiaxed grain structure. A n y twins are not observed in the S Z . Grains o f the SZ
46
·
Materials Processing and Texture
Effect of T e x t u r e o n F r a c t u r e Limit S t r a i n in Friction Stir P r o c e s s e d A Z 3 1 Β M g Alloy
are significantly refined b y recrystallization during F S P , and average grain size of the SZ was 3.6 μηι. In the T Z , the m i x e d grain structure consisting of fine and coarse grains is observed.
20 Till •!>«: 3 d * KotttUMl
1«
Trawl wood: 11 mm/t
ε
•
m
a. I
12
Β
I
ε υ
4
Trt«H*p«*
0
1.7 1.8 1.* Plunge depth ( m m )
a
•
2
1
*
3
4
5
20 HUNTSFCPTK l.IRAE Tih mile; 3 dc« ?m«l 12 n a / i
ifmi
ϊ's *
'S
ο * 500
Ί
16
ΐ
12
Λ.
2
Η
\Xw\mlt
T i l t angle (deg)
2«
1"
Pluftac depth 1 λ mm tUMOonal t p w d 3 0 0 RPM
^
w
η E
»
2 υ
4
1000 1500 2000 2500 JO00
2
4
6
8
10
12
14
T r a v e l speed (mm/s)
Rotational speed ( r p m )
Figure 5. Effect of F S P parameters on grain size in the stir z o n e .
<0
0.45 0.40 0.35 0.30 0.25 0.20 0.15 0.10 0.05 0.00
•
Ο
0
2
4
6
FSPed material BM
8 10 12 14 16 18 20 22 Grain size (μηι)
Figure 6. Effect of grain size o n fracture limit major strain of the stir z o n e . Effect of F S P p a r a m e t e r s o n grain size in the stir zone of the F S P e d alloy is s h o w n in Fig. 5. The grain size hardly d e p e n d s o n the plunge depth and the tilt angle of the F S P tool, while it
Materials Processing and Texture
·
47
Effect of T e x t u r e o n F r a c t u r e Limit S t r a i n in Friction Stir P r o c e s s e d A Z 3 1 Β M g Alloy
increases with increasing rotational speed remarkably and with decreasing travel speed slightly. T h e r e w a s n o large variation in grain size inside the stir zone o f all FSPed alloys. In general, the grain size in the stir zone is related with the heat input during friction stirring, and the higher heat input causes the large grains in the stir z o n e . Since it is k n o w n that the heat input increases w i t h increasing rotational speed and with decreasing travel s p e e d , the larger grains are produced in the stir zone at both t h e higher rotation and the slower travel of t h e F S P tool. Effect o f grain size o n fracture limit £^ ,, in the stir zone is presented in F i g . 6. T h e grain size increases with grain size up to 10 μπι, b e y o n d which it decreases. T h i s tendency is different w h i c h s h o w e d the higher ductility in the finer grained from those of the previous papers, " M g alloys. T h i s result m e a n s that the other microstructural factors affect the fracture limit ε^αmore strongly t h a n the grain size in the stir z o n e . Since the m e c h a n i c a l properties of M g alloy are significantly dependent o n the crystallographic t e x t u r e , the texture analysis w a s c o n d u c t e d in the stir zones. 15
o;t
2
4 1 2 1 J
11
Figure 7. (0002) pole figures of the B M and the stir zones p r o d u c e d using Parameters 1, 8. 9, 10 and 1 1 . (0002) pole figures o f the B M and the C C regions of the SZs produced using Parameters 1, 8, 9, 10 and 11 a r e s h o w n in Fig. 7. F r o m this figure, effect o f rotational speed o n texture in the stir zone can be seen. T h e B M exhibits a typical hot-rolled texture c o m p o n e n t s h o w i n g that the (0002) basal plane is parallel to t h e plate surface. T h e C C regions of the S Z s also h a v e the strong accumulation of the (0002) basal plane near the N D . but the accumulated poles of the (0002) basal plane incline from the N D to the F S P D . T h e inclined angle varied b e t w e e n 5 and 35
48
·
Materials Processing and Texture
Effect of T e x t u r e o n F r a c t u r e Limit S t r a i n in Friction Stir P r o c e s s e d A Z 3 1 Β M g Alloy
d e g , but n o clear d e p e n d e n c e o f the FSP parameters o n the inclined angle w a s found in this study.
Figure 8. Distribution of the (0002) pole figure in the stir zone produced using Parameter 2. Texture distribution inside the stir zone p r o d u c e d using P a r a m e t e r 2 is g i v e n in Fig. 8. All regions in the stir zone exhibit the similar texture c o m p o n e n t s s h o w i n g that the (0002) basal planes are strongly accumulated near the N D , although the (0002) poles of all regions do not align to the N D exactly. This result m e a n s that t h e texture distribution is roughly h o m o g e n e o u s in this stir z o n e . T h e similar feature of the texture distribution w a s found in all stir zones produced in this study. Previous s t u d i e s ' ' have examined the texture distribution in the stir zones of A l and M g alloys by electron backscattered diffraction ( E B S D ) m e t h o d . M o s t of t h e m have s h o w n that the shear texture c o m p o n e n t is d o m i n a n t in the stir zone, and that the shear plane is roughly parallel to the pin c o l u m n surface of the tool, w h i c h leads to rotating distribution o f shear texture ' F r o m these results, it is p r o p o s e d that the material around t h e pin surface in the stir z o n e . m o v e m e n t mainly arises from the shear deformation a l o n g the pin surface during friction stirring. In M g alloys, P a r k et a l . have s h o w n that the material flow during friction stirring mainly activates t h e (0002) basal plane slip systems, w h i c h p r o d u c e s the strong basal plane textures along the pin surface in the stir zone. O n the other hand, this study obtained the roughly h o m o g e n e o u s texture distribution in the stir zone of M g alloy after F S P . The (0002) basal plane w a s roughly parallel to the plate surface in all regions o f the stir z o n e . T h e s e results imply that the material m o v e m e n t along the shoulder surface is d o m i n a n t for the texture formation in M g alloys during F S P . Jin et a l . reported that shear deformation along t h e shoulder surface w a s d o m i n a n t during F S W of 2 m m thick 5000-series Al alloy w h e n the tool having the shoulder diameter o f 15 m m w a s used. This 4 1 6
1 7
1 4 , 1 6
1 7
14
18
Materials Processing a n d Texture
·
49
Effect of T e x t u r e o n F r a c t u r e Limit Strain in Friction Stir P r o c e s s e d A Z 3 1 Β M g Alloy
result w o u l d be due t o sweeping t h e r e m n a n t of shear deformation along t h e p i n surface b y shear along the shoulder surface, because the material m o v e m e n t along the shoulder surface continues for a while after the pin passes. T h e present study also used the F S P tool h a v i n g t h e shoulder diameter of 15 m m for t h e 2 m m - t h i c k M g plate, so that the shear along the shoulder surface would be a reason for the h o m o g e n e o u s texture distribution in the stir zone.
Figure 9. Definition of the angle (0) b e t w e e n the direction perpendicular to the tensile direction and the basal plane n o r m a l .
0.40| 0.35 0.30 0.25 0.20 0.151 0.10 0.051 a nn '""θ
1
' · . ;.·,
•
• Ο
5
10
15
20
FSPed material I BM I
25
30
35
Figure 10. Effect of angle φ on fracture limit major strain of the stir z o n e . All FSPed alloys produced in this study had the similar texture in the stir zones. T h e (0002) basal planes w e r e a c c u m u l a t e d a r o u n d the N D , but t h e (0002) poles inclined 5 to 35 deg from the N D . K o i k e et a l . e x a m i n e d the texture d e p e n d e n c e o n the active slip systems in rolled AZ61 M g alloy with the basal plane texture during tensile test. T h e y s h o w e d that the tensile elongation w a s related to the tilt angle b e t w e e n n o r m a l direction of the basal plane and the tensile direction. W h e n the tilt angle w a s small (the normal direction of the basal plane inclined toward the tensile direction largely), basal < a > slip w a s a d o m i n a n t deformation m e c h a n i s m and 1 9
50
·
Materials Processing and Texture
Effect of T e x t u r e o n F r a c t u r e Limit S t r a i n in Friction Stir P r o c e s s e d A Z 3 1 Β M g Alloy
a small elongation w a s obtained. Otherwise, prismatic slip w a s activated in M g alloy at the tilt angle close to 90 d e g , leading to good tensile ductility. A c c o r d i n g to this finding, the angle (φ) b e t w e e n the direction perpendicular to the tensile direction (the N D in this study) and the basal plane n o r m a l w a s estimated from the (0002) pole figures of the C C regions, and then effect in the stir zone w a s plotted in Fig. 9. T h e fracture limit e^or of φ on fracture limit £ j decreases with increasing φ, except for the B M , w h i c h is in g o o d agreement with the previous r e s u l t . This result suggests that the fracture limit ε„φ is affected by d e g r e e of the activity of obtained at the lower φ by m o r e non-basal slip systems and that t h e larger fracture limit ε„φ activity of non-basal slip systems. T h e lower fracture limit e j of the B M m a y be due to the large spread of the (0002) pole, i.e. the relatively r a n d o m texture. F r o m Fig. 6, it m a y seem that the grain size hardly affects the fracture limit ε„α/ , but since the present study produced only the 0 larger than 20 deg in all stir zones having the smaller grain size than 10 μπι, effect of grain size o n ductility could not be properly discussed in this study. H o w e v e r , m o s t of the previous studies h a v e reported the better ductility in the finer grained M g alloy, so that the M g alloy with the finer grains should h a v e the better fracture limit ε„φ, at the same φ. This study s h o w s that the crystallographic texture is also a major microstructural factor governing the fracture limit ε^φτ in the F S P e d M g alloy as well as the grain size, and suggests that both the fine grains and the strong basal plane texture parallel to t h e plate surface need to be achieved during F S P for production of the highly formable M g alloy. m a
0 r
19
Γ
Γ
ma 0r
ΟΓ
CONCLUSION T h e present study e x a m i n e d effect of microstructure, especially grain size and texture, on fracture limit major strain under the uniaxial tension in the stir zone of the F S P e d A Z 3 1 B M g alloy. All stir zones exhibited the strong accumulation of (0001) basal plane near the plate surface, w h o s e formation w a s attributed to shear deformation along the tool shoulder during FSP. T h e fracture limit major strain increased with increasing the grain size u p to 10 μπι, beyond w h i c h it decreased. O n the other hand, the fracture limit major strain w a s higher w h e n the angle between n o r m a l direction of the basal plane and the tensile direction w a s closer to 90 d e g in the stir zones. T h i s study suggests that both the fine grain structure and the strong basal plane texture parallel to the plate surface are microstructural features to improve the formability in the FSPed M g alloy. ACKNOWLEDGEMENT T h e authors are grateful to M r . A . H o n d a for technical assistance and a c k n o w l e d g e Prof. J. K o i k e , Prof. K. Ikeda, Prof. H . Sato and Prof. Z.J. W a n g for their helpful discussion. O s a w a Scientific Studies Grants Foundation is gratefully a c k n o w l e d g e d . A part of this study w a s financially supported by the Japanese Ministry of Education, Science, Sports and Culture with a Grant-in-Aid from the Global C O E p r o g r a m of the Materials Integration International Center for Education and Research at T o h o k u University. REFERENCES Ά . Couret and D. Caillard, In Situ Study on Prismatic Glide in M a g n e s i u m - I. The Rate Controlling M e c h a n i s m , Acta Metall., 3 3 , 1447-54 (1985). A . Yamashita, Z. Horita, and T.G. L a n g d o n , Improving the Mechanical Properties of M a g n e s i u m and a M a g n e s i u m Alloy through Severe Plastic Deformation, Mater. Sei. Eng. A, A300, 142-47(2001). 2
Materials Processing and Texture
·
51
Effect of T e x t u r e o n F r a c t u r e Limit S t r a i n in Friction Stir P r o c e s s e d A Z 3 1 Β M g Alloy
3
T . M u k a i , M . Y a m a n o i , H. W a t a n a b e , and K. Higashi, Ductility E n h a n c e m e n t in A Z 3 1 M a g n e s i u m Alloy by Controlling Its Grain Structure, Scripta Mater., 4 5 , 89-94 ( 2 0 0 1 ) . J . Koike, T. K o b a y a s h i , T. M u k a i , H. W a t a n a b e , M. Suzuki, K. M a r u y a m a , and K. Higashi, T h e Activity of N o n - B a s a l Slip S y s t e m s a n d D y n a m i c Recovery at R o o m T e m p e r a t u r e in F i n e Grained A Z 3 1 B M a g n e s i u m Alloys, Acta Mater., 5 1 , 2 0 5 5 - 6 5 ( 2 0 0 3 ) .
4
5
R . S . Mishra, M . W . M a h o n e y , S.X. M c F a d d e n , N . A . M a r a , and A . K . Mukherjee, H i g h Strain Rate Superplasticity in a Friction Stir Processed 7075 Al Alloy, Scripta Mater., 4 2 , 163-68 (2000). 6
P . B . Berbon, W.H. Bingel, R . S . Mishra, C C . B a m p t o n , and M . W . M a h o n e y , Friction Stir Processing: A T o o l to H o m o g e n i z e N a n o c o m p o s i t e A l u m i n u m Alloys, Scripta Mater., 4 4 , 6 1 - 6 6 (2001). 7
M . W . M a h o n e y , C G . R h o d e s , J.G. Flintoff, R.A. Spurling, a n d W . H . Bingel, Properties o f Friction-Stir-Welded 7 0 7 5 T 6 5 1 A l u m i n u m , Metall. Mater. Trans. A, 29 A, 1955-64 (1998). Y . S . Sato, H. K o k a w a , M . E n o m o t o , and S. Jogan, Microstructural Evolution o f 6063 A l u m i n u m During Friction-Stir W e l d i n g , Metall. Mater. Trans. A, 3 0 A , 2 4 2 9 - 3 7 (1999). K . V . Jata and S.L. Semiatin, C o n t i n u o u s D y n a m i c Recrystallization D u r i n g Friction Stir W e l d i n g o f H i g h Strength A l u m i n u m A l l o y s , Scripta Mater., 4 3 , 743-49 (2000). J . - Q . Su, T.W. N e l s o n , R. Mishra, and M . M a h o n e y , Microstructural Investigation Friction Stir W e l d e d 7 0 5 0 - T 6 5 1 A l u m i n u m , Acta Mater., 5 1 , 713-29 (2003). 8
9
I0
" S . H . C . Park, Y . S . Friction Stir W e l d of Y . S . Sato, S.H.C. Highly F o r m a b l e M g I 2
Sato, a n d H. K o k a w a , Effect o f M i c r o - T e x t u r e on Fracture Location in M g Alloy A Z 6 1 During Tensile Test, Scripta Mater., 4 9 , 161-66 (2003). Park, A . M a t s u n a g a , A . H o n d a , a n d H . K o k a w a , N o v e l P r o d u c t i o n for Alloy Plate, J. Mater. Sei., 4 0 , 6 3 7 - 4 2 (2005).
I 3
Y . S . Sato, A . Sasaki, A. S u g i m o t o , A. H o n d a , and H. K o k a w a , E n h a n c e m e n t o f Formability in M a g n e s i u m Alloy A Z 3 1 B via Friction Stir Processing, Mater. Sei. Forum, 5 3 9 - 5 4 3 , 3 7 7 5 - 3 7 8 0 (2007). S . H . C Park, Y . S . Sato, and H. K o k a w a , Basal Plane T e x t u r e and F l o w Pattern in Friction Stir W e l d of a M a g n e s i u m Alloy, Metall. Mater. Trans. A, 3 4 A , 9 8 7 - 9 4 (2003). Y . S . Sato, M . Urata, H. K o k a w a , and K. Ikeda, Hall-Petch Relationship in Friction Stir W e l d s of Equal Channel Angular-Pressed A l u m i n i u m Alloys, Mater. Sei. Eng. A, A 3 5 4 , 2 9 8 - 3 0 5 (2003). Y . S . Sato, H. K o k a w a , K. Ikeda, M. E n o m o t o , S. Jogan, and T. H a s h i m o t o , Microtexture in the Friction-Stir-Weld o f an A l u m i n u m Alloy, Metall. Mater. Trans. A, 3 2 A , 941-48 ( 2 0 0 1 ) . D . P . Field, T.W. N e l s o n , Y . H o v a n s k i , a n d K.V. Jata, Heterogeneity o f Crystallographic Texture in Friction Stir W e l d s of A l u m i n u m , Metall. Mater. Trans. A, 3 2 A , 2 8 6 9 - 7 7 (2001). H . Jin, S. Saimoto, M . Ball, and P.L. Threadgill, Characterisation o f Microstructure and Texture in Friction Stir W e l d e d Joints o f 5 7 5 4 and 5182 A l u m i n u m Alloy Sheets, Mater. Sei. Technol.,\l, 1605-14(2001). I 4
15
I 6
, 7
18
19
J . K o i k e and R. O h y a m a , Geometrical Criterion for the Activation of Prismatic Slip in A Z 6 1 M g Alloy Sheets Deformed at R o o m T e m p e r a t u r e , Acta Mater., 5 3 , 1 9 6 3 - 7 2 (2005).
52
·
Materials Processing and Texture
M I C R O S T R U C T U R E A N D T E X T U R E ANALYSIS OF FRICTION STIR WELDS OF COPPER Saukkonen, T., Savolainen, K., M o n o n e n , J., and H ä n n i n e n , H. Helsinki University of Technology, Laboratory of Engineering Materials E s p o o , Finland ABSTRACT C o p p e r is readily j o i n e d by friction stir welding ( F S W ) due to its relatively low melting point. The stacking fault energy (SFE) of copper is comparatively low, resulting in a discontinuous d y n a m i c recrystallization ( D R X ) m e c h a n i s m operating at the temperature and strain rate range typical for F S W . T o study the textures in the weld area, several F S W e d copper ( C u - O F E P ) samples w e r e studied. A s a m p l e representing the o p t i m u m process conditions (tool temperature 850 °C) of a full scale 50 m m thick F S W w e l d of copper capsule for spent nuclear fuel w a s obtained from Swedish N u c l e a r Fuel and W a s t e M a n a g e m e n t C o ( S K B ) . Smaller sized w e l d s w i t h varying processing p a r a m e t e r s w e r e a l s o prepared t o study t h e effects of the processing parameters deviating from optimal. T h e microstructure w a s studied with optical microscopy, F E G S E M and E B S D . N e w metallographical m e t h o d s w e r e developed to be able to distinguish the different zones in t h e w e l d area. T h e texture in F S W w e l d s of copper is weak and it varies slightly according to the location in the weld, because the direction of the flow of material in F S W varies constantly in the weld area. T h e m a i n texture c o m p o n e n t found in F S W w e l d s of copper is < 1 0 1 > fibre texture. INTRODUCTION Friction stir w e l d i n g ( F S W ) is a solid state w e l d i n g p r o c e s s d e v e l o p e d b y T W I and it is mainly used to j o i n a l u m i n i u m alloys. T h e basic m e t h o d is relatively simple. T h e rotating tool consisting of a p i n and a shoulder are plunged into the material until the shoulder contacts the surface. T h e friction b e t w e e n the tool and the workpiece causes t h e heating of the material typically u p to 0.6 - 0.9 T of the material and softens it e n o u g h for the tool to be able to travel through the material. T h e m o v e m e n t of the plasticized material in connection and around the pin allow the solid joint to be formed. T h e melting point of p u r e copper is 1083 °C, thus, possible temperatures of the material in connection with the tool are 6 0 0 - 9 0 0 °C. T h e actual temperature is controlled mainly by the rotation and traverse speeds as well as the plunge d e p t h and tilt angle o f t h e tool. O n t h e material side t h e temperature and other p r o c e s s parameters are the result o f the flow stress d e p e n d i n g on the actual strain rate - temperature c o m b i n a t i o n s and the stacking fault energy (SFE) of the material. C o p p e r is a F C C metal and can b e defined to have a m e d i u m l o w S F E (78 m J / m ) . A s c o m p a r e d to a l u m i n i u m with high SFE (typically 166 m J / m ) , copper behaves differently in F S W and the k n o w l e d g e gathered in a large n u m b e r of studies of the F S W of a l u m i n i u m cannot b e directly applied t o copper. It h a s been established that the m e c h a n i s m operating in F S W of copper is dynamic . recrystallisation ( D R X ) and the texture developed in D R X of copper is found to be w e a k The m e c h a n i s m of d y n a m i c recrystallisation is different dependent o n the S F E of the material. W h e n the S F E is high, as in a l u m i n i u m , the dynamic recrystallisation is termed as the continuous dynamic recrystallisation ( C D R X ) . T h e high S F E a l l o w s the d y n a m i c recovery ( D R V ) to operate. The high dislocation density caused by deformation in F S W is arranged as small angle boundaries. With increased strain energy the misorientation of the small angle boundaries m
2
2
1
1,2,3
53
T h e M i c r o s t r u c t u r e a n d T e x t u r e A n a l y s i s of Friction Stir W e l d s of C o p p e r
increases continuously until they are transformed to high angle b o u n d a r i e s and n e w grains are formed. T h u s , in C D R X n e w grains are formed from the existing grains, w h i c h are divided into smaller grains. T h e orientation and the texture of the n e w grains are linked to the texture o f the parent grains. O n the contrary, in the l o w or m e d i u m S F E materials the planarity o f the slip hinders the cross slip and c l i m b of the dislocations causing the activation energy for D R V to be higher than that for the nucleation of n e w grains. N e w grains g r o w and continuing deformation causes n e w strain energy gradients in the g r o w i n g grains until they are h i g h e n o u g h again for a n e w grain to nucleate. This m e c h a n i s m of d y n a m i c recrystallisation is called discontinuous dynamic recrystallisation ( D D R X ) . T h e m e c h a n i s m of D R X , especially in connection of a h i g h strain rate process such as F S W , is not very well understood. The necklace structure typical of some D R X processes can b e explained by bulging of the grain b o u n d a r i e s . The initiation o f D R X in copper in connection to triple j u n c t i o n s h a s b e e n f o u n d . T h e nucleation of n e w grains and the extensive t w i n n i n g typical t o D R X in l o w S F E materials have b e e n stated to explain the randomness o f the texture as a result of D R X . 4
5
The resulting texture and the grain size in the steady state of D R X are independent o f the grain structure o f the parent material. In Figure 1 a b o u n d a r y area b e t w e e n the weld and the base material of F S W e d o x y g e n free copper, can b e seen. T h e base material, seen in the right h a n d
Figure 1. T h e Euler m a p s and the pole figures near the b o u n d a r y o f dynamically recrystallized zone. In the upper part of the image the s a m e area is seen as a continuous orientation image. T h e twin boundaries are highlighted with light lines. T h e area o f the separate i m a g e s in the middle o f the picture is 3 6 6 x 2 8 8 μπι. side o f the image,is in a cold-rolled state with very large grain size. Only a few grains are covered in the image, but the details o f the typical rolling texture of F C C materials can be detected in the corresponding pole figure. On the left hand side of the i m a g e a typical equiaxed grain structure of F S W e d copper as a result of D R X c a n be seen. T h e corresponding pole figure indicates almost r a n d o m texture. Because of the heat generated b y F S W , a n a r r o w area of the
54
·
Materials Processing and Texture
T h e M i c r o s t r u c t u r e a n d T e x t u r e A n a l y s i s of Friction Stir W e l d s of C o p p e r
base material is recrystallized b y static recrystallization ( S R X ) a s can b e seen b y the typical annealing texture in the pole figures in the m i d d l e of the i m a g e . W h e n the b a s e material is in hot rolled state with m u c h smaller grain size, t h e resulting grain structure in the w e l d area is similar to the w e l d n u g g e t in Figure 1 and it w o u l d be v e r y difficult to distinguish the b o u n d a r y between the weld and the base material. EXPERIMENTAL T h e full scale spent nuclear fuel copper canister F S W w e l d i n g of the lid and the capsule w a s performed in S K B canister laboratory in S w e d e n using the temperature controlled s y s t e m . T h e temperature m e a s u r e d from the inside of the tool w a s 850 °C. T h e thickness of the copper canister wall is 50 m m . Small laboratory scale (pin length a b o u t 5.3 m m ) w e l d s ' with varying process parameters w e r e prepared to study the effect process parameters o n the quality of the w e l d s . T h e cross-sections o f t h e w e l d s w e r e prepared w i t h typical s a m p l e preparation m e t h o d s and various n e w m e t h o d s w e r e studied to b e able to discern different zones of the weld area. Samples used for E B S D and channeling contrast studies w e r e polished electrolytically as a final step. T h e S E M and E B S D studies w e r e performed w i t h Zeiss Ultra 55 F E G S E M . E B S D acquisition and analyses w e r e performed with the Oxford Instruments N o r d l y s II c a m e r a and Channel 5 software. T h e material used is a special grade of oxygen-free copper with t h e addition of about 4 0 p p m of p h o s p h o r o u s ( C u - O F E P ) , specially designed for copper canister purposes. 6
1
8
MICROSTRUCTURAL STUDIES M o s t of t h e results c o n c e r n i n g t h e laboratory scale samples are presented elsewhere . F S W w a s first applied to j o i n i n g of aluminium alloys and it is still today the main application area. T h u s , the t e r m s and definitions describing the w e l d metallography w e r e developed based on the metallographical details typical to the cross-section o f F S W e d aluminium, e.g. the onion rings describing the n u g g e t area are usually m e n t i o n e d as typical features of F S W . Actually, in o u r e x p e r i m e n t s the onion rings w e r e only found, w h e n the welded surfaces w e r e intentionally left uncleaned and the welding p a r a m e t e r s were c h o s e n t o p r o d u c e too cold conditions. T h u s , the onion ring like features detected o n t h e cross-section of c o p p e r F S W w e l d indicates a nonsatisfactory quality of the weld. Figure 2 b s h o w s typical cross-section of w e l d area of F S W e d copper o f a spent nuclear fuel canister lid . 8
4
Figure 2 . F S W w e l d between the lid and the copper capsule in S K B canister construction, a) location o f the w e l d and b) typical metallurgical cross-section of the w e l d , in w h i c h the weld area is approximately m a r k e d b y d a s h e d lines. T h e thickness o f the w e l d is 50 m m . 4
3
Materials Processing and Texture
·
55
T h e M i c r o s t r u c t u r e a n d T e x t u r e A n a l y s i s of Friction Stir W e l d s of C o p p e r
A s can be seen in F i g u r e 2, t h e actual F S W w e l d area c a n n o t b e easily detected using conventional optical metallography. T o study the texture and other characteristics of the w e l d , n e w m e t h o d s w e r e d e v e l o p e d to be able t o detect the different z o n e s of the weld area. Figure 3 s h o w s a specially etched and t h e n electropolished cross-section of a S K B copper canister w e l d revealing the different processing zones of the weld.
Figure 3 . M a c r o s c o p e i m a g e of the l o w e r part of the cross-section o f a S K B copper canister F S W weld using the w h i t e light side illumination to reveal the interesting zones of the weld area. E v e n if the welding position is horizontal, see Figure 1, the images o f the weld area are presented in the vertical direction in this report. T h e areas used by E B S D for texture m e a s u r e m e n t s are marked b y letters. S E M i m a g e in Figure 4 s h o w s the origin o f the sharp contrast differences in Figure 3 . T h e specific grooved m a r k s in the left h a n d side o f Figure 4, i.e. in the b a s e material are formed by the heavy lateral flow of the electrolyte causing some kind of tails or g r o o v e s connected to C u S inclusions s h o w n in detail in the enlargement in the lower left h a n d side corner of Figure 4. T h e solubility of S in C u increases rapidly at temperatures o v e r 8 0 0 °C. Obviously, high temperature and large deformation d e g r e e typical t o F S W process cause the C u S inclusions t o dissolve rapidly in the weld area and in cooling to precipitate in a m u c h finer form s h o w n in the right hand side of Figure 4. T h e fine precipitation of C u S in F S W area m a y b e the explanation for the limited grain growth in F S W e d area at high temperatures, see F i g u r e 5. F i g u r e 6 s h o w s E B S D m a p s of the microstructure in t h e nugget area. 2
2
2
56
·
Materials Processing and Texture
T h e M i c r o s t r u c t u r e a n d T e x t u r e A n a l y s i s of Friction Stir W e l d s of C o p p e r
Figure 4. S E M i m a g e s h o w i n g the origin o f sharp contrast in the area m a r k e d b y C in Figure 3.
Figure 5. S E M i m a g e of the b o u n d a r y area between the w e l d and the base material after h y d r o g e n annealing for 3 0 min at 850 ° C s h o w i n g notable grain growth in t h e base material c o m p a r e d that in the F S W w e l d area.
Materials Processing and Texture
57
T h e M i c r o s t r u c t u r e a n d T e x t u r e A n a l y s i s of Friction Stir W e l d s of C o p p e r
Figure 6. E B S D m a p s of t h e n u g g e t area, a) pattern quality m a p w i t h grain boundaries as black lines and b) Euler m a p with twin boundaries marked by light lines. T h e a m o u n t o f twin boundaries is about 2 8 % . T h e length of the scale bar in the i m a g e s is 2 0 0 μηι. TEXTURE STUDIES Pole figures w e r e calculated using the stereographic projection, upper h e m i s p h e r e with half width of 10°. T h e O D F s w e r e calculated with the series expansion m e t h o d with Gaussian half with of 5°. Figure 7 s h o w s the pole figures of the area m a r k e d b y A and Β in Figure 3 corresponding to the copper lid and capsule materials, respectively. It c a n be seen that the texture of the lid and the capsule material differs noticeably. A l s o the grain sizes are different.
58
·
Materials Processing and Texture
T h e M i c r o s t r u c t u r e a n d T e x t u r e A n a l y s i s of Friction Stir W e l d s of C o p p e r
Figure 9. Pole figures of a cold laboratory scale weld with a texture of 4.5 times r a n d o m and O D F sections of φ 2 = 4 5 " and 75° s h o w i n g <101 > fibre texture c o m p o n e n t s . Figure 8 s h o w s the pole figures in different locations m a r k e d in Figure 3. Texture varies slightly in the different areas possibly due to different processing conditions in different locations. The direction of the flow o f material varies constantly in the w e l d area. A l s o the relatively small n u m b e r of grains covered c a u s e variability in texture. W h e n trying to detect possible texture differences in small areas as in A r e a C in Figure 4 , large n u m b e r of grains cannot be covered. In m a n y locations, especially in the m i d d l e areas of the weld, features s h o w i n g < 1 0 1 > fibre texture c o m p o n e n t s are present. T h e textures in the upper part of the cross-section of the weld not shown in Figure 1, near the surface o f the weld, were similar to the texture in A r e a F ( w e a k textures with varying degree of features of < 1 0 1 > fibre texture). H a s e g a w a et a l found that texture in D R X of copper is dependent on Ζ value ( Z e n e r - H o l l o m o n parameter) being sharper in high Ζ conditions (low temperature, high strain rate increases Z ) . Figure 9 s h o w s the texture in laboratory scale s a m p l e welded in h i g h e r Ζ conditions. T h e texture is sharper and clearly concentrated to < 1 0 1 > fibre texture. It can be concluded that the < 1 0 1 > fibre texture is the prominent c o m p o n e n t in the rather w e a k a n d varying texture o f F S W e d copper. 5
CONCLUSIONS T h e friction stir welding of copper is a c o m p l e x process involving the high strain rate material deformation and flow as well as dissolution of certain inclusions, e.g. C u S , into solid solution of c o p p e r affecting the grain structure a n d the grain size. In welding of a full scale copper canister the grain structure in F S W e d copper j o i n t s is formed by D R X consisting of equiaxed grains with only a w e a k texture in the whole weld area w h e n welded with o p t i m u m processing parameters using t h e temperature controlled process. T h e resulting grain structure and 2
Materials Processing and Texture
59
T h e M i c r o s t r u c t u r e a n d T e x t u r e A n a l y s i s of Friction Stir W e l d s of C o p p e r
the grain size are dependent o n the process parameters, not on the grain size or texture of the base material. T h e small laboratory scale w e l d s w e r e used to study the effect of process parameters deviating from the optimal conditions to the quality and texture of the F S W weld. It w a s found that the w e a k texture sharpened w h e n the temperature of the F S W p r o c e s s w a s lowered (higher Z) and the texture concentrated to the orientation of < 1 0 1 > fibre texture. Texture m e a s u r e m e n t using E B S D can be used to evaluate the w e l d i n g p a r a m e t e r s for finding the o p t i m u m process w i n d o w for F S W . It can readily reveal the subtle c h a n c e s in the grain structure, w h i c h are not possible to detect by conventional metallographic techniques. Indications o f increasing texture components m a y be the result of nonoptimised process p a r a m e t e r s and/or insufficient cleaning of the weld surfaces resulting in oxides or other particles in the processing zones. ACKNOWLEDGEMENTS The authors w i s h to t h a n k T e k e s (Finnish Funding A g e n c y for T e c h n o l o g y and Innovation) and K Y T 2 0 1 0 (Finnish Research P r o g r a m m e on Nuclear W a s t e M a n a g e m e n t ) project for funding. Copper F S W e l d s provided by the Swedish N u c l e a r Fuel and W a s t e M a n a g e m e n t C o ( S K B ) and Lars Cederqvist ( S K B ) w e r e essential for the study. REFERENCES 1 Savolainen, K., M o n o n e n , J., Saukkonen, T., H ä n n i n e n , H., and Koivula, J. "Friction Stir Weldability of C o p p e r A l l o y s " , 5th International Friction Stir Welding Conference, Metz, France (2004). 2 M . H a s e g a w a a, M . Y a m a m o t o b , H. Fukutomi, Formation M e c h a n i s m of Texture During D y n a m i c Recrystallization in γ-TiAl, Nickel and C o p p e r E x a m i n e d by Microstructure Observation and Grain B o u n d a r y Analysis Based o n L o c a l Orientation M e a s u r e m e n t s , Acta Materialia 51 3 9 3 9 - 3 9 5 0 (2003). 3 D. T. M c D o n a l d , F. J. H u m p h r e y s and P. S. Bate, N u c l e a t i o n and Texture D e v e l o p m e n t During D y n a m i c Recrystallization of Copper, Materials Science Forum, 4 9 5 - 4 9 7 , 1195-1200 (2005). 4 H u m p h r e y s , F.J. and Hatherly, M . "Recrystallization and Related A n n e a l i n g Processes", 2 n d ed, Oxford, U K : Elsevier Science Ltd, 4 9 7 p , I S B N 0-08-041884-8 (1995) 5 H. Miura, T. Sakai, H. Hamaji and J. J. Jonas, "Preferential Nucleation of D y n a m i c Recrystallization at Triple j u n c t i o n s " , Scripta Materialia, 50 (1), 65-69 (2004). 6 Cederqvist, L. " F S W to Seal 50 m m Thick C o p p e r Canisters - A W e l d that Lasts for 100,000 Y e a r s " , 5th International Friction Stir Welding Symposium, M e t z , France (2004). 7 Swedish N u c l e a r Fuel and W a s t e M a n a g e m e n t C o , " K a p s e l för A n v ä n t Kärnbränsle, Svetsning vid Tillverkning o c h Förslutning", report R - 0 6 - 0 4 (2006). 8 Savolainen, K., S a u k k o n e n , T., M o n o n e n , J., and H ä n n i n e n , H., "Entrapped Oxide Particles in C o p p e r Friction Stir W e l d s " , 7th International Symposium on Friction Stir Welding, Awaji Island, Japan (2008), to be published.
60
Materials Processing and Texture
Texture and Anisotropy in Steels
This Page Intentionally Left Blank
THE E F F E C T S O F T E X T U R E ON T H E CORROSION R E S I S T A N C E O F SS304 STAINLESS STEEL M. Azzi and J. A. Szpunar Department of Materials Engineering, McGill University, Montréal, Canada E-mails: marwan.azzi(g)mail.mcgill,ca. [email protected]
ABSTRACT Stainless steels are well k n o w n for their high corrosion resistance properties. In this paper, the effect of surface texture on the corrosion resistance properties of SS304 stainless steels was investigated. Different surface textures were produced for this study, namely {001 }<100>, {110}<110> and {111}<110>. Electrochemical Impedance Spectroscopy (EIS) was performed t o characterize the general corrosion resistance of the stainless steel surfaces. Potentiodynamic polarization was performed to characterize the resistance to localized corrosion (Pitting) of the stainless steel surfaces. The poles figures were measured using x-ray diffraction technique and the orientation distribution functions (ODFs) were calculated using Resmat software. The results showed that surfaces with (111) preferred planes have the highest resistance to general corrosion and pitting. The resistance to general corrosion of the surfaces with (111) preferred planes was 258 Μ Ω . α η compared to 219 Μ Ω . α η and 82 Μ Ω . α η for the surfaces with (001) (respectively (110)) preferred planes. In addition, the pitting potential was found to be 0.46V for (111) textured surfaces compared to 0.29V and 0.15V for (001) (respectively (110)) textured surfaces. It has been shown that the corrosion resistance increased with increasing planar atomic density of the surface. 2
2
2
1-INTRODUCTION Stainless steels (Fe-Cr alloys) are widely used in industry due to their good resistance to corrosion. This is mainly due to the very thin (few nanometer) oxide layer, known as passive layer, that protect the surface against dissolution. The passive layer is composed of Cr-oxide and Fe-oxide, the first being the responsible for the excellent corrosion properties of the material [1]. SS304 is the most versatile and most widely used stainless steel, available in a wider range of products and forms. It is known as " 1 8 / 8 " which means 1 8 % of Chromium and 8% of Nickel in iron. Pitting corrosion, one of the most severe and destructive types of localized attack, is one of the major problems related to the stability of stainless steels used in various applications [2]. An extensive research has been done to improve both the general and pitting resistance of stainless steels. Ion implantation of SS304 by molybdenum and titanium has been shown to improve the corrosion resistance [3]. Plasma nitriding, which is used to increase the wear resistance of stainless steels, has been also used to increase the corrosion resistance. Shahryari and al. [4] have successfully improved the pitting resistance of biomedical grade 316LVM stainless steel by cyclic potentiodynamic polarization in sodium nitrate. In this paper, the role of surface texture in the corrosion resistance properties of SS304 stainless steels was investigated. Electrochemical Impedance Spectroscopy and Potentiodynamic polarization were performed to characterize the corrosion resistance of the stainless steel surfaces.
63
T h e Effects of T e x t u r e o n t h e C o r r o s i o n R e s i s t a n c e of S S 3 0 4 S t a i n l e s s S t e e l
2- E X P E R I M E N T A L 2-1 Materials Hot rolled AISI304 stainless steel sheets of 5 m m thickness were used as raw materials in order to prepare samples with different textures. The sheet surface had initially a strong cubic texture {001}<100>. Using different angles for cutting as shown in Fig. 1. three different types of surface were produced; Type 1 : Samples whose surface normal coincides with z-axis, as shown in Fig. 1 a. Type 2: Samples whose surface normal coincides with vector (1,1,0), as shown in Fig. l b . Type 3: Samples whose surface normal coincides with vector W ( l , l , - 1 ) , as shown in Fig. l c . Texture measurement was performed for the three different types of surface, because the texture was inhomogeneous through thickness. The rolling direction for each type of surface is shown in Fig. 1.
Fig. 1 Preparation of samples with different textures; (a) Type 1 (b) Type 2 (c) Type 3 For corrosion experiments, specimens with 7mmx5mm rectangular area were prepared. Copper wires were connected to the back side of the specimens using silver conducting epoxy to ensure proper electrical contact for electrochemical measurements. The specimens were then mounted in Tecknovit 4004 resin and polished down to 1200 grade silicon carbide paper. NaCl 1% was used as electrolyte. 2-2 Corrosion test A standard three-electrode cell configuration was used for all corrosion experiments. The stainless steel specimens served as working electrodes, the counter electrode was a graphite rod and the reference electrode was a Saturated Calomel Electrode (SCE). The three electrodes were connected to an Autolab PGSTAT302 potentiostat equipped with a Frequency Response Analyser for EIS measurements. Corrosion experiments consisted of electrode stabilization in the NaCl 1% solution for 1 hour at Open Circuit Potential (OCP). Next, the Electrochemical Impedance Spectroscopy (EIS) was performed in the frequency range between 50 K H z and 10 m H z with a voltage perturbation of 10 m V . Finally, the potentiodynamic test was performed starting from 50 m V below the OCP, in the cathodic zone, to a potential at which a current density of 1 m A / c m was reached. The potential scan rate was 1 mV/sec. The corrosion current was found using the Tafel interpolation technique. 2
64
Materials Processing and Texture
T h e Effects of T e x t u r e o n t h e Corrosion Resistance of S S 3 0 4 Stainless Steel
EIS technique allows us to investigate the general corrosion resistance of an electrode in corrosive medium. Briefly, this technique consists on generating a low-amplitude sine wave of voltage, with different frequencies, between the working electrode and the reference electrode, and measuring the current between the working electrode and the counter electrode. At each frequency, the impedance is given by Z{w) = ψ \ =
-MxpOwO
=
^ e x p O V ) = Z expO«>) = Z cos«>) + Z sin(<*>) 0
0
0
Where E is the voltage amplitude imposed, in V I is the current amplitude measured, in A Z is the impedance magnitude w is the frequency, in rd/sec t is the time, in seconds φ is the phase shift between the imposed voltage signal and the measured current signal In this paper, the impedance spectra were presented in Bode diagrams where the impedance is plotted with the log frequency on the x-axis and the absolute values of the impedance |Z| = Z and the phase shift φ on the y-axis. The potentiodynamic polarization allows us to characterize the surface in terms of resistance to localized corrosion (known as pitting). The pitting resistance is characterized in terms of pitting potential, Epi , which is defined as 0
0
0
0
t
Epit
Ebreak * E rr CO
is the corrosion potential, and Ebreak is the breakdown potential which corresponds to the Where E potential at which there is a rapid increase in the current due to the breakdown of the passive layer. c o l r
3- R E S U L T S A N D DISCUSSION 3-1 Texture measurement Texture measurements were performed for the three different types of stainless steel surfaces. The Orientation distribution function for the surface of type 1 (respectively type 2 and type 3) is presented in Fig. 2 (respectively Fig. 3 and Fig. 4 ) . The results showed that type 1 surface had {001 }<100> preferred orientation, type 2 surface had {110}<110> preferred orientation, and type 3 surface had {111}<110> preferred orientation. Since the direction of the plane on the surface is not important for corrosion resistance, only the preferred plane will be taken into consideration for the rest of the paper; the type 1 surface (respectively type 2 and type 3) will be referred as (001) preferred plane surface (respectively (110) and (111) preferred plane surfaces).
Materials Processing a n d Texture
·
65
T h e Effects of T e x t u r e o n t h e C o r r o s i o n R e s i s t a n c e of S S 3 0 4 S t a i n l e s s S t e e l
Fig. 2: Orientation distribution function of AISI304 stainless steel sample of typel
Fig. 3 : Orientation distribution function of AISI304 stainless steel sample of type2
66
·
Materials Processing and Texture
T h e Effects of T e x t u r e o n t h e C o r r o s i o n R e s i s t a n c e of S S 3 0 4 S t a i n l e s s S t e e l
Max=
ο ^ ,
v
-
/
r
y ·
λ
Wo 9©
10.7
—
1.00
—
2.90
—
4.90
—
6.80
—
8.70
) (@) ι
·* φι =0°,90°
ψ
^ /
ο
ο
Φ » 0 , 90 r Fig. 4: Orientation distribution function of AISI304 stainless steel sample of type3
3-2 EIS results The EIS spectra of SS304 samples with (001), (111) and (110) preferred planes are presented in Bode diagram in Fig. 5. The left y-axis is the absolute value of the impedance Log(Zo), and the right yaxis is the phase shift φ. It is clearly seen that the impedance of the SS304 (110) was much lower than the impedances of the other two samples, which were close. The difference between (111) and (001) samples w a s only noticeable on the Arg(Z) curve at low frequencies; the (111) sample showed slightly higher phase shift. The impedance spectra were interpreted in terms of appropriate Equivalent Electrical Circuits (ECC). Since the impedance spectra were characterized by two-time-constant behavior (two peaks on the Arg(Z) plot), it was necessary for the fitting procedure to use an ECC with two time constant elements, such as capacitor. The E C C are shown in the inset to Fig. 6. It describes the behavior of a process characterized by the high-frequency time constant (Qj, Ri), and the low-frequency time constant (Qj, R2). According to Bou-Saleh and al. [2], the physical meaning of the H F time constant can be associated with faster charging/discharging and charge transfer processes, respectively represented by the constant phase element, Q j , and resistance to charge transfer, Ri, in parallel combination. However, the LF time constant can be related to slower mass transport processes in the oxide phase. R represents the solution resistance of the test electrolyte between the working electrode and the reference electrode. The fitting procedure showed that a better agreement between the experimental and theoretical spectra was obtained when pure capacitance was replaced by the constant phase element, Q. its impedance is expressed as s
2
Where Y„ is the parameter of Q (Fsec""'cm" ), w is the angular frequency and η is the empirical exponent of Q that always lies between 0.5 and 1 [4]. If n = l , Q can be related to pure capacitance, if n
Materials Processing and Texture
67
T h e Effects of T e x t u r e o n t h e C o r r o s i o n R e s i s t a n c e of S S 3 0 4 S t a i n l e s s S t e e l
Taking into account the physical meaning of the E C C . the sum of the two resistances. Ri + R , is related to the surface's total resistance to general corrosion. 2
Fig. 5 EIS spectra of SS304 samples with (001), (111), and (110) preferred planes presented in Bode plot Fig. 6 shows the EIS spectrum and the fitting curve of SS304 (111) sample, as well as the E C C used for the fitting procedure. The agreement between the experimental (triangles) and the modeled (solid line) data was excellent, thus justifying the use of the proposed ECC. The data obtained from the simulation of all EIS spectra are presented in Table I. ( I l l ) surface was found to have the highest total corrosion resistance ( R i + R ) , followed by (001) surface and then (110) surface. It should b e mentioned that the SS304 stainless steel is an austenitic stainless steel with f e e structure, and therefore (111) plane is the most densely packed plane and (110) has the lowest planar density compared to the other two planes. This indicated that it is possible to correlate between the planar atomic density and the corrosion resistance of a surface. These results showed that it is more difficult to dissolve metallic ions from a surface that has a higher atomic density than a surface with lower atomic density. 2
68
Materials Processing and Texture
T h e E f f e c t s of T e x t u r e o n t h e C o r r o s i o n R e s i s t a n c e of S S 3 0 4 S t a i n l e s s S t e e l
Fig. 6 EIS spectrum and fitting curve of the SS304 sample with (111) preferred plane as well as the electrical circuit used for simulation
Table I The data obtained from Sample R (Ω) SS304 ( 1 Π ) 4.9 SS304(001) 5.2 SS304(110) 4.5 s
the EIS spectra simulation Qi (Y.) (F) 0.72X10" 0.76X100.53x10"
QiOO 4
4
0.90 0.90 0.87
Ri (Ω) 25.2x10 14.95x10 1.44xl0 J
3
Qi (Y.) (F) 0.20x10" 0.27Χ10" 0.69x10"
Q2
(«)
4
3
4
4
0.89 0.84 0.71
233x10 204x10 80x10
3
3
3
3-3 Pitting corrosion resistance The polarization curves of SS304 surfaces with (001), (111), and (110) preferred planes are presented in Fig. 7. Table II presents the pitting potentials and corrosion rates obtained from the polarization curves of the different samples. For (111) and (001) surfaces, the onset of pitting was very clear on the polarization curves, while for (110) surface, it was difficult to locate the beginning of pitting since the current was continuously increasing with the potential. This indicated that the pitting susceptibility of (110) surface was relatively high compared to the other surfaces and that the pits started to forms at the early anodic potentials. ( I l l ) surface was found to have the highest pitting potential and the lowest corrosion rate, indicating that the most densely packed plane in the stainless steel structure had the highest corrosion resistance. These results confirmed the results obtained in the EIS technique.
Materials Processing and Texture
·
69
T h e Effects of T e x t u r e o n t h e C o r r o s i o n R e s i s t a n c e of S S 3 0 4 S t a i n l e s s S t e e l
Fig. 7 Polarization curves of SS304 surfaces with (111). (001) and (110) preferred planes Table II Corrosion rates and the pitting potentials obtained from the polarization curves Sample Corrosion rate (A/cm ) Pitting potential (V) SS304(111)
1
SS304(001)
2x10'"
SS304(110)
I
x
1
0
.
6x10'"
039 1
Approx.0.15
4- S U M M A R Y It was found that the surface texture plays an important role in the corrosion properties of SS304 stainless steel. A (111) preferred plane surface exhibited a significant improvement in pitting resistance, whereas a (110) preferred plane surface decreased the pitting resistance of SS304. The pitting potential was 0.46 V for the (111) oriented surface, 0.29 V for the (001) oriented surface and 0.15V approximately for the (110) oriented surface. In addition, the resistance to general corrosion of the stainless steel surfaces was improved with (111) preferred planes. This difference in corrosion properties is due to the difference in the planar atomic density the different planes have; a higher planar atomic density resulted in a better corrosion resistance. REFERENCES [1] E.E. Stansbury and R.A. Buchanan, Fundamental of electrochemical corrosion, A S M International, 2000. [2] Z. Bou-Saleh, A. Shahryari, S. Omanovic, Thin Solid Film 515 (2007) 4727-4737 [3] X. B. Tian, Z. M . Zeng, T. Zhang, Β. Y. Tang and P. K. Chu, Thin solid film 366 (2000) 150-154 [4] A. Shahryari, S. Omanovic. J. A. Szpunar, Materials Science & Engineering C, Article in press [5] H.G. Kim, S. Ahn, J. Kim, S.J. Park, K.R. Lee. Thin Solid Film 475 (2005) 291-297
70
·
Materials Processing and Texture
M I C R O S T R U C T U R A L C H A R A C T E R I Z A T I O N O F DUAL PHASE STEELS BY M E A N S OF ELECTRON M I C R O S C O P Y 2
Liesbeth Barbé' and Kim V e r b e k e n * ' C R M (Centre for Research in Metallurgy) Technologiepark 903c B-9052 Zwijnaarde Belgium [email protected] d e p a r t m e n t of Metallurgy and Materials Science Faculty of Engineering Ghent University (UGent) Technologiepark 903 B-9052 Ghent Belgium * Postdoctoral fellow with the Fund'for Scientific Research-Flanders (Belgium) (F.W.O. Vlaanderen) ABSTRACT Dual phase (DP) steels have a low yield strength due to the soft ferrite matrix and a high tensile strength due to the hard dispersed martensite second phase islands. This makes them excellent steels for use in light weight automobile structural parts. A profound understanding of microstructureproperties relationships of advanced high strength steels is of crucial importance to the rational development of these steels. This demands the ability to perform advanced quantification of the complex multiphase microstructure, addressing the full range of microstructural properties that are relevant to the macroscopic mechanical behaviour. The ferrite-martensite microstructure will be studied by means of Light Optical Microscopy, Scanning Electron Microscopy and Electron BackScatter Diffraction. The martensite fraction in ferritemartensite DP steels will be measured both on a FEG-SEM and a LaBö-SEM. The influence of the microscope resolution on the determination of the martensite fraction by means of the Image Quality (IQ) factor will be investigated. INTRODUCTION Improved passenger comfort and safety features have resulted in the increase of the weight of recent automotive models, despite the efforts to reduce weight, fuel consumption and emissions. A considerable research effort has therefore focused on the development of high strength sheet steels for the automotive industry. The use of high strength, thin gauge sheet steel offers a reasonable solution of weight reduction. A n increased crash-worthiness to improve passengers safety must, however, also b e achieved. High strength steels offer excellent properties for use in impact energy absorbing structural parts such as crumple zones at the front and backside of cars, as they gradually cushion the shock. In addition, the passenger compartment must transform into a stiff, rigid frame with anti-intrusion steel reinforcements embedded in the car doors, to minimise the passenger compartment collapse in case of severe impact. Furthermore, the high strength steels offer more car body styling freedom and lead to a large variety of exterior appearances as a result of excellent formability properties. Dual phase (DP) steels consist of a soft ferrite matrix, with fine dispersed martensite islands. It combines a low yield stress due to the soft ferrite matrix with a high tensile strength as a result of the hard second phase. The amount of martensite mainly determines the strength of the steel and typically varies between 5 and 5 0 % . The steel obtains its specific properties from the low temperature transformation of intercritical austenite to martensite. This transformation results in a non-uniform
71
M i c r o s t r u c t u r a l C h a r a c t e r i z a t i o n of D u a l P h a s e S t e e l s b y M e a n s of E l e c t r o n M i c r o s c o p y
dislocation distribution and internal stresses in the material. The austenite, enriched in carbon during the processing, may not transform entirely into martensite, resulting in the presence of small amounts of retained austenite. The amount of retained austenite depends on the hardenability of the transforming austenite phase and the size of the austenite particles. For most industrially processed dual phase steels, the presence of small amounts of other phases, such as bainite and, to a lesser extent, pearlite, can be tolerated. They do not significantly change the specific dual phase properties as long as a sufficient amount of low-temperature martensite has been formed. Important for enhancing the mechanical performance of these kinds of high strength steels and the ability to make production processes more efficient is the availability of a set of tools to quantify their microstructural properties as a first step towards facilitating an appreciation of the correlation between processing conditions and mechanical properties. These tools should comprise a suite of nonsubjective metallographic procedures providing efficient, robust, reproducible and accurate microstructural quantification. The sub-micron size of the different microstructural components in high strength steels like DP steels renders traditional methods for microstructural quantification to a large extent ineffective as the resolution of conventional light optical microscopy is far too low. Electron BackScatter Diffraction (EBSD) has been found to be a powerful technique for studying the local crystallographic properties of a number of microstructures'' . 2
EXPERIMENTAL The research is performed on an industrial DP600 cold rolled steel grade. The light optical micrographs were taken on a ZEISS optical microscope equipped with a pixel-based K O N T R O N K S 4 0 0 * image analysis software program. After mechanical polishing, the samples were either etched with a 2 % Nital solution or a LePera etchant . For the SEM study, a Philips X L 3 0 ESEM, equipped with a LaB -filament, was used. The acceleration voltage was 25 kV, the spot size 3.5-4.5 and the working distance (WD) around 10. The samples were lightly etched with Beraha to reveal the microstructural components. For the EBSD measurements a working distance (WD) of 20 was used. The polished and unetched specimen was tilted over 75° and the diffraction patterns were captured by a digital C C D camera. The data were analysed with the TSL Orientation Imaging Microscopy software. 3
6
RESULTS A N D DISCUSSION A Nital etching was used to reveal the grain boundaries of the ferrite phase, while the martensite phase appeared dark (Figure 1). Pictures were taken in three directions in the plate, i.e. R D T D plane, R D - N D plane and N D - T D plane. An automatic measurement of the ferrite grain size was only possible in the RD-TD plane. The ferrite grain size was ~ 1 0 p m and the ferrite grains had a stretch factor of 1.12. The ferrite grains seen on the two other planes have a compressed and elongated appearance according to the rolling process. "Squeezed" martensite bands could be seen in between the ferrite grains.
72
Materials Processing and Texture
M i c r o s t r u c t u r a l C h a r a c t e r i z a t i o n of D u a l P h a s e S t e e l s b y M e a n s of E l e c t r o n M i c r o s c o p y
Figure 1. Overview of the L O M of the DP600 steel, scale bar ΙΟμπι. A LePera etching was used to determine the amount of martensite which appeared white or dark brown depending on the plane. The martensite volume fraction varied between 8% measured on the R D - T D plane and 1 5 % on the N D - T D plane. The LePera etching was also able to reveal the ferrite grain boundaries and a ferrite grain size of 5.4μπι was measured on both the R D - N D and ND-TD plane. The ferrite grains had a stretch factor of 1.31. With L O M , grains smaller than l - 2 u m cannot be determined and measured correctly. This has to be taken into account because almost 1 5 % of the grains are thought to be in this range. In order to quantify the martensite volume fraction in a more accurate way, the samples etched with Beraha, were investigated by S E M where a clear contrast between the ferrite and martensite phase was visible (Figure 2). Quantitative Image Analysis (QIA) performed on 6 photographs of the RD-ND plane at magnification 2000 lead to an average martensite volume fraction of 2 0 % . This value is somewhat higher as the maximum value of 1 5 % obtained with L O M but can be explained by the fact that with SEM small martensite islands are recognised and measured as well. Figure 2 also shows that a minor volume fraction of bainite is present in the material which can be seen as lamellae or less smooth islands or clusters of small particles.
Materials Processing and Texture
·
73
M i c r o s t r u c t u r a l C h a r a c t e r i z a t i o n of D u a l P h a s e S t e e l s b y M e a n s of E l e c t r o n M i c r o s c o p y
Figure 2. S E M of the DP600 steel, Beraha etchant. 2
The unetched samples were also investigated by E B S D . An area of 180*380μπι was scanned with a step size of 0.5μπι on the R D - N D plane. A clean up of the data was performed by means of "neighbour CI correlation" with a minimum CI of 0.1 and "grain dilation" with a minimum of 3 pixels. Separation of the martensite and the ferrite phase could be done by using the Image Quality (IQ), which was noticeably lower for the martensite since the crystallographic structure of this phase is more distorted which leads to less clear diffraction patterns. As will be discussed below, an IQ of 45 appeared to be a good measure to distinguish between martensite and ferrite. The I Q and Inverse Pole Figure (IPF) maps are given in Figure 3.
Figure 3. IQ map (left) and IPF map for the polygonal ferrite (middle) and the martensite phase (right). As can be seen on the IQ map, the banded structure of martensite islands, which was also found with L O M (cf. Figure 1), can easily be identified. It was calculated that the fraction of martensite phase
74
·
Materials Processing and Texture
M i c r o s t r u c t u r a l C h a r a c t e r i z a t i o n of D u a l P h a s e S t e e l s b y M e a n s of E l e c t r o n M i c r o s c o p y
was about 1 3 % . When this amount is compared with previously mentioned fractions of martensite, a possible underestimation of the martensite fraction can be recognized. This can probably be attributed to the step size used, 0.5 μπι, which will result in loss of the smallest martensite islands. The choice of a step size is, however, a compromise between the scanning of a statistically relevant area in a reasonable time and the size of the second phase particles. Due to the limited resolution of the LaB(,filament equipment, it was chosen to repeat the present measurement with a smaller step size on a high resolution microscope as is shown below. Another important parameter in the quantification process is the image quality factor. This parameter describes the quality of an electron backscatter diffraction pattern. While the IQ is dependent on the material and its condition, it is not an absolute value but a function of the technique and parameters used to index the pattern as well as other factors such as the video processing. The IQ parameter is the sum of the detected peaks in the Hough transform. Thus, any parameters that affect the operation of the Hough transform would influence the resulting IQ parameter calculation. In addition, changing conditions in the microscope or video processing (such as simply changing the contrast and brightness) will also affect the IQ values. Similarly, the orientation independence is introduced into the calculation due to the fact that different bands of varying intensity will appear depending on the orientation . It has to be stressed that the cut-off value of 4 5 , used in this work, is arbitrary chosen. If a too low cut-off value is chosen, one risks, on the one hand, to exclude the interiors of the larger martensite blocks which show nice diffraction patterns and, on the other hand, to include a too high fraction of grain boundaries. Indeed, grain boundaries will show a low IQ because of the fact that the low band contrast in these regions can be related with the distorted crystal lattices. W u et al. have developed a method to characterise complex microstructures based on the different image quality of each particular phase. This method would avoid that one has to choose a cutoff value for the image quality. The starting point of this new approach is the fact that the different types of ferrite produced at different temperatures during cooling, such as polygonal, acicular and bainitic ferrite, will have different densities of dislocations or sub-grain boundaries. The effects of image processing on the IQ values are minimised through a normalisation procedure. The multi-peak model is based on the fact that an IQ distribution of one micro-constituent is a symmetric peak and thus the whole IQ distribution curve can be divided into several normal distributions, each of which represents a certain micro-constituent. A normalized IQ distribution was calculated and fitted by means of Gaussian distributions (Figure 4a). According to Wu et al. the grain boundary regions (GBR) will result in a peak at low IQ values. When these G B R points are subtracted from the IQ distribution, two Gaussian curves can be fitted to the IQ curve (Figure 4b). According the L O M (Figure 1), the DP600 steel contains only two phases, namely martensite and ferrite and hence those two Gaussian peaks can be related to the martensite and ferrite phase. When the peak surfaces are calculated, 3 8 % martensite is found which is significantly higher than the value obtained from the LOM and SEM pictures. Although the method proposed by W u et al. has proven to be valuable for D P steels with a high fraction of martensite , the present work indicates that it is not easily applicable to a DP steel with lower martensite fractions that was studied in this work. 4
5
5
5
Materials Processing and Texture
·
75
Microstructural C h a r a c t e r i z a t i o n of D u a l P h a s e S t e e l s b y M e a n s of E l e c t r o n M i c r o s c o p y
Nom.llMd IQ
N«™ita* U
Figure 4. Normalized IQ distribution curves for the DP600 steel (a) without and (b)with grain boundary region correction. The same sample was measured on a F E G - S E M in order to investigate the influence of the resolution of the microscope on the possibility to differentiate between martensite and ferrite. The step size was 0.15μπι and an area of 57*71 μηι was measured. The data was cleaned up with a "grain dilation" procedure with minimum grain size 5. When the IQ map (Figure 5) is compared with the one of the LaBj measurement (Figure 3). it is clear that more details of the martensite phase are visible. For the present case, image quality values lower than 60 were considered to correspond with martensite. This resulted in a martensite volume fraction of 8% which is lower than the one for the LaBn filament measurement. It was found that defining a cut-off value for the high resolution measurement was even more difficult because the high resolution caused the band contrast in the martensite phase to be better and consequently the average I Q of martensitic data points was higher. Consequently, it was even more difficult to distinguish between grain boundaries and martensite. Another reason for the low martensite fraction might also be the fact that a relatively small area needed to be taken because the lower step size would otherwise increase too much the measurement time. 2
Figure 5. IQ of the D P 6 0 0 steel measured on a FEG-SEM. Besides the quantification of the phases is also the crystallographic texture an important feature during the microstructural characterization of the material. The E B S D technique is well suited to determine
76
Materials Processing and Texture
M i c r o s t r u c t u r a l C h a r a c t e r i z a t i o n of D u a l P h a s e S t e e l s b y M e a n s of E l e c t r o n M i c r o s c o p y
separately the texture of both phases present. Figure 6 shows the
Figure 6. Texture of the ferrite (left) and martensite (right) in the DP600 steel. CONCLUSIONS An industrial DP600 steel was investigated by means of light optical microscopy, scanning electron microscopy and electron backscatter diffraction. A LePera etching can reveal information on the ferrite phase such as grain size and grain shape factor and will reveal the martensite phase. This method has its limitations as the martensite islands are very small, Ιμηι and less. A SEM investigation on a Beraha etched samples was able to measure the martensite volume fraction in an accurate way. The disadvantage of this technique is the loss of all information on the ferrite grains. Finally, the material was investigated by E B S D , both on a SEM with L a B filament and a FEG-SEM. The distinction between martensite and ferrite can be made by the image quality factor. This method, however, is more subjective as the analysis with normal SEM but gives more information on both phases such as the texture. 6
ACKNOWLEDGEMENTS This work was carried out with a financial grant from the Research Fund for Coal and Steel of the European Community. FOOTNOTE * The grain shape factor is determined by fitting an ellipse to the points making up a grain. One measure of the shape is the stretch factor. The stretch factor is defined as the length of the major axis divided by the length of the minor axis. REFERENCES ' V. Randle, Microtexture determination and its applications, 1982, London, The institute of metals. D.J. Dingley and V. Randle, Microtexture Determination by Electron Back-Scatter Diffraction,./. Mai. Science, 2 7 , 4 5 4 5 - 6 6 (1992). F.S. LePera, Improved Etching Technique to Emphasize Martensite and Bainite in Hing-Strength Dual-Phase Steel, Journal of Metals, 32, 38-9 (1980). T S L * O I M analysis for Windows, 2005. J. Wu, P.J. Wray, C.I. Garcia, M . Hua and A.J. DeArdo, "Image quality analysis: A new method of characterizing microstructures", ISIJ International, Vol. 4 5 , No. 2, 2005, p p . 254-262 2
3
4
5
Materials Processing and Texture
77
This Page Intentionally Left Blank
G R A I N S I Z E E F F E C T O N T H E T E N S I L E B E H A V I O R O F A F e M n C T W I P S T E E L IN RELATION WITH THE MICROSTRUCTURE A N D TEXTURE EVOLUTIONS 1,2
3
David B a r b i e r , Nathalie G e y ' , Alain Sébastien , Michel H u m b e r t
1
1
Laboratoire d ' E t u d e des Textures et Application aux Matériaux ( L E T A M ) , U M R C N R S 7 0 7 8 , Université d e M e t z , 5 7 0 0 0 Metz, France Laboratoire de Physique et M é c a n i q u e des Matériaux ( L P M M ) , U M R C N R S 7 5 5 4 , Université de M e t z , 5 7 0 0 0 M e t z , France Arcelor Research, V o i e R o m a i n e , B P 30320, 5 7 2 8 3 Maizières-les-Metz
2
3
ABSTRACT T h e influence of the grain size in a high m a n g a n e s e austenitic T W I P (Twinning Induced Plasticity) steel h a s b e e n investigated o n samples deformed in tension. T h e twinning activity varies with the grain size and influences the mechanical behavior. In the earlier stage of deformation, mechanical twinning is favored in the small grain-sized steel, resulting in a higher hardening rate. For h i g h deformation levels, the large grain-sized steel exhibits a higher twin v o l u m e fraction. This reduces its texture sharpness and m a i n t a i n s a high strain hardening rate. INTRODUCTION T h e deformation of F e - M n - C T W I P steels involves t w o strain m e c h a n i s m s : dislocation glide and twinning. This leads to outstanding mechanical properties, c o m b i n i n g high resistance and high ductility. T h e v o l u m e fraction of t w i n s increases with the strain, reducing continually the m e a n free path of the dislocations. This is called the T W I P effect (Twinning Induced Plasticity), w h i c h c a n be interpreted as a dynamical Hall-Petch effect . In recent years, the deformation behavior o f fee materials with different Stacking Fault Energies ( S F E ) h a v e b e e n characterized. T h e s e studies mainly c o n c e r n M P 3 5 N , α - B r a s s or Hadfield s t e e l s , w h i c h also h a v e h i g h m a n g a n e s e content. A four stage strain hardening has often b e e n reported, independent from the material S F E and grain size. This behavior has been analyzed in relation to the microstructure evolution and in particular the twin activation. The characteristic microstructural events h a v e b e e n discussed o n the basis of optical and Transmission Electron M i c r o s c o p y analysis that are either limited in resolution or suffer from a lack of representativeness. T h e influence of t w i n n i n g o n the strain hardening rate h a s widely been analyzed. Different m o d e l s h a v e been proposed to take the d y n a m i c Hall-Petch effect into account. H o w e v e r , the influence of the grain size o n the t w i n activity and the corresponding strain hardening rate is not well understood. This subject still raises a lot of attention and interest. M o r e o v e r less studies report the influence of the texture on the twinning activation and consequently o n the m e c h a n i c a l behavior. In this contribution, the tensile behavior of high m a n g a n e s e austenitic T W I P steels has b e e n investigated in relation to microstructure and texture evolutions. T w o grain-sized grades have b e e n analyzed and the grain size effect o n t h e microstructure and texture evolution has been discussed. E B S D m a p p i n g s in a F E G S E M w e r e performed to analyze the twinning microstructure with a high resolution and a m o r e statistical aspect. 1-2
3
4
1
5
79
G r a i n S i z e Effect o n T e n s i l e B e h a v i o r of a F e M n C T W I P S t e e l
MATERIAL AND EXPERIMENTAL PROCEDURE Material The investigated material is a n austenitic F e - 2 2 % M n 0 . 6 % C (weight % ) steel, provided by ArcelorMittal Maizières Research S A (ArcelorMittal G r o u p ) . A c c o r d i n g to its chemical composition, mechanical twinning w a s expected to occur under a deformation at room temperature. Industrial slabs w e r e first hot rolled to a thickness of 3 m m and then 5 0 % reduced by cold rolling. Recrystallization and annealing treatments w e r e adapted to obtain different grain-sized s h e e t s . The present study focuses o n the 2.6 μηι (small grain S G ) and the 15.5μπι (large grain L G ) grades. Their initial textures w e r e rather similar with a texture index of 1.23. 6
Mechanical testing Tensile tests w e r e performed at r o o m temperature in the Rolling Direction ( R D ) at a strain rate of 1 0 ' V . The true stress and inelastic strain w e r e calculated as well as the normalized strain
hardening
ratei — ) / G y de )
(the
0
material
shear m o d u l u s
temperature). Go w a s determined according to the relation G °
G
0
g
=62000
MPa
at
room
(the Y o u n g m o d u l u s Ε
2(1 + ν )
and the Poisson coefficient υ w e r e m e a s u r e d experimentally). To calculate the strain hardening rate, the experimental stress-strain curve w a s fitted with a 7 degree polynomial function. Texture and microstructural characterization T h e deformation textures w e r e determined from the experimental { 2 0 0 } , {220} and {111} poles figures (PFs). T h e s e pole figures w e r e m e a s u r e d by X - R a y Diffraction on a S i e m e n s D 5 0 0 0 goniometer using the Cr Κ radiation. The experimental data w e r e processed to obtain the Orientation Distribution Function f(g) ( O D F ) . T h e orthorhombic symmetry w a s observed on the experimental pole figures and w a s i m p o s e d for t h e O D F calculation. T h e sharpness o f t h e texture 2 w a s evaluated with the Texture Index (TI) (given by Tl = jf(g) dg ). T h e v o l u m e fractions (VF) α
7
of specific texture c o m p o n e n t s w e r e calculated by taking into account a misorientation of 10° around the strict orientations (overlapping of the considered texture c o m p o n e n t s w a s avoided). High resolution orientation m a p s w e r e acquired in a 6 5 0 0 F J E O L Field Emission G u n Scanning Electron M i c r o s c o p e ( F E G - S E M ) equipped with the Channel 5 E B S D system ( H K L Technology). S a m p l e s were prepared by mechanical and ionic polishing using a B a l T e c R E S 0 1 0 system. T h e microstructure and the local texture evolutions w e r e analyzed using the Kikuchi Pattern Quality ( K P Q ) m a p s and the R D - I n v e r s e Pole Figure (IPF) m a p s . RESULTS AND DISCUSSION M e c h a n i c a l behavior The true stress and the normalized strain hardening rate are plotted as a function of the strain in figure 1. C u r v e s in b l u e correspond to g r a d e S G a n d those in red t o g r a d e L G (the mechanical behavior of intermediate grain-sized grades is reported in reference 6). For both grades, a combination of high strength and high ductility w a s observed, mainly due to a remarkable high hardening rate up to large strains. In fact, the mechanical twins act as obstacles for dislocation gliding ( d y n a m i c Hall P e t c h effect) ' and promote this high hardening rate. 8
80
Materials Processing and Texture
9
G r a i n S i z e Effect o n T e n s i l e B e h a v i o r of a F e M n C T W I P S t e e l
Figure 1. T r u e stress-strain curves and normalized strain h a r d e n i n g for the SG grade (in blue) and for the L G (in red). M o r e o v e r , the strain h a r d e n i n g evolves according to different regimes. In particular, after a strong d e c r e a s e (stage A ) , the strain hardening rate increases (stage B ) for both grades. A fairly similar evolution h a s been reported by A S G A R I and E L - D A N A F o n other low S F E materials. T h e first increase o f the strain h a r d e n i n g curve (stage B ) w a s linked to the occurrence o f the first mechanical t w i n s ' . 3
3
4
4
The austenitic grain size affects the m e c h a n i c a l behavior by t w o w a y s : (i) T h e flow stress increases with decreasing grain size (the increase is o f 1 5 0 M P a w h e n the grain size decreases from 15.5μηι to 2.6μηι). This effect is classically described by the HallP e t c h relation (1) o =o y
0
+ kd ^ ( 1 )
w h e r e Oy is the yield stress, σο and k two material parameters (respectively 132.15 M P a and 4 4 8 . 6 M P a / μ η ι " ) . (ii) T h e grain size also alters the strain h a r d e n i n g rate. In fact, the transition from o n e strain h a r d e n i n g stage to the next h a p p e n s at l o w e r strain levels for grade S G . In particular, stage Β (increase o f strain h a r d e n i n g ) starts at ε =0.02 for grade S G a n d at ε,=0.10 for grade L G . H o w e v e r , the stress level for the onset o f stage Β is close for both grades (about 5 3 0 M P a ) . A c c o r d i n g to , it c a n be a s s u m e d that twinning o c c u r s earlier for grade SG than for grade L G . This point will be further discussed latter in the text by m e a n s o f microstructural observations. 2 5
(
4
C o n s e q u e n t l y , for l o w strains, a h i g h e r strain h a r d e n i n g level is observed for grade S G c o m p a r e d to grade L G . F o r h i g h e r strains, the t e n d e n c y is inverted with a small advantage for grade L G u p to the end o f the deformation. Texture evolution T h e deformation textures at 8 = 0 . 5 2 are reported for both grades in figure 2 with the {100}/{110}/{111} pole figures a n d the O D F section at q>2=45°. The tensile direction w a s parallel t o the prior sheet rolling direction (corresponding to the vertical direction o f the PFs). T h e main orientations observed on the O D F sections are indicated in table 1 b y m e a n s of the {hkl} notation a n d t h e Euler A n g l e s according to the B u n g e convention. The corresponding s y m b o l s are reported o n t h e O D F section at p h i 2 = 4 5 ° . T h e v o l u m e fraction of each c o m p o n e n t is given in figure 3 . T h e highest S c h m i d F a c t o r (SF) for slip and twinning (considering t h e {111}<110> slip s y s t e m s and the {111}<112> twinning systems) h a v e been calculated - see table 1. The m a c r o s c o p i c stress w a s considered equal to the microscopic o n e . t
Materials Processing a n d Texture
·
81
G r a i n S i z e Effect o n T e n s i l e B e h a v i o r of a F e M n C T W I P S t e e l
Table 1 : main crystallographic orientations, their orientations in the φ2=45° section and the maximal value o f the associated Schmid factor.
T h e tensile deformation produces a sharp austenitic texture. A s seen on the P F s , this texture is characterized by a p r o n o u n c e d < 1 1 1 > / / R D fiber and a less pronounced < 1 0 0 > / / R D fiber (with reinforcement around specific orientation c o m p o n e n t s ) . T h e s e t w o fibers are often observed in austenitic s t e e l s ' and are induced by the crystallographic rotations resulting from dislocation gliding (as confirmed by Taylor m o d e l i n g " ) . O f course, in case of T W I P steels, twinning influences the slip activity . In particular, twinning can modify the relative intensity of each f i b e r " . Moreover, when the twin v o l u m e fraction b e c o m e s sufficient, the twinning orientations h a v e t o b e taken into account t o describe the deformation texture. 0
82
·
Materials Processing and Texture
G r a i n S i z e Effect o n T e n s i l e B e h a v i o r of a F e M n C T W I P S t e e l
E v e n if b o t h g r a d e s exhibit t h e s a m e m a i n c o m p o n e n t s (Rotated Brass and Copper), several differences occur according to the grain size. Indeed, the texture index is higher for grade SG ( T I = 4 . 0 6 ) t h a n for g r a d e L G ( T l ' ° = 2 . 3 4 ) . T h i s increase o f texture sharpness from grade LG t o S G , results from a n increase of orientation densities a r o u n d specific c o m p o n e n t s (see the densities in figure 2 and the corresponding v o l u m e fractions in figure 3 - in particular for the Rotated B r a s s and t h e C o p p e r c o m p o n e n t s belonging to t h e < 1 1 1 > / / R D fiber). T h u s , the decrease of the grain size m a y p r o m o t e the grain rotations t o w a r d the deformation c o m p o n e n t s . SG
F i g u r e 3 . V o l u m e fractions of the different texture c o m p o n e n t s according to the grain size ( S G in blue and L G in red). H o w e v e r , the role o f twins should also be taken into account w h e n analyzing the grain size effect o n deformation textures o f T W I P steels. Indeed the v o l u m e fraction of twin orientations seems higher for grade L G than for grade SG. In particular, the Copper Twin orientation that is obtained by mechanical twinning from the C o p p e r orientation , is more pronounced for grade L G t h a n for grade S G (see figure 2 and 3). T h u s it could also be assumed that the v o l u m e fraction of t w i n s is higher for grade L G than for grade SG. This is in agreement with previous works reporting that the v o l u m e fraction of twins increases with increasing grain s i z e ' . Further discussion about the role of twinning on the deformation texture requires additional microstructural characterizations in relation with the texture evolution. 3
Microstructure evolution T h e microstructure evolution and the twinning activity have been studied with the E B S D technique. Figure 4 presents a characteristic b a n d contrast m a p and RD-1PF m a p at ^ = 0 . 5 2 for each grade. A c c o r d i n g to the color of the standard triangle, grains in red belong to the < 1 0 0 > / / R D fiber ( C u b e / G o s s / C o p p e r t w i n orientations) and grains in blue belong to the < 1 1 1 > / / R D fiber (Rotated Brass/Copper). In regions w h e r e the orientations could not be determined, the grey color is given by the band contrast.
Materials Processing and Texture
·
83
G r a i n S i z e Effect o n T e n s i l e B e h a v i o r of a F e M n C T W I P S t e e l
Figure 4: E B S D analyses ( K P Q + I P F ) of b o t h grades ( S G o n the left / L G o n the right side) tensile strained at ε = 0 . 5 2 The b a n d contrast m a p reveals the c o m p l e x microstructure in w h i c h t h e t w i n s are organized in stacks, or b u n d l e s , of n a n o - t w i n s . In case of grade S G , the grain size varies in a large interval from about 0.7 to 15μπι. Grains larger than 5μπι are less n u m e r o u s ( 1 0 % in number) but cover 4 0 % o f the material surface. O n the band contrast m a p , one c a n see that t h e twins mainly form in the larger grains w h e r e a s grains with a size lower than 1.5μηι in majority exhibit no twins. T h i s confirms a grain size effect on the t w i n n i n g activation. 2 1 3
A strong d e p e n d e n c e b e t w e e n grain orientation and t w i n n i n g activity is also observed. T h e grains in blue, belonging to the Rotated Brass and C o p p e r orientations exhibit intense twinning activation. T h e grains in red, belonging to the G o s s o r C u b e orientations, are only deformed by slip. In fact, the < 1 0 0 > / / R D fiber has an orientation a d v a n t a g e for slip ( S F = 0 . 4 1 ; S F = 0 . 2 3 ) . For the < 1 1 1 > / / R D fiber, slip and twinning are m o r e competitive ( S F = 0 . 3 1 ; S F = 0 . 2 7 ) . T h u s the d e v e l o p m e n t of a pronounced < 1 1 1 > / / R D fiber p r o m o t e s t w i n n i n g in the polycrystal. T h e v o l u m e fraction of grains belonging to this fiber increases progressively during the deformation. sllp
lwto
l w m
sllp
T w o main differences can be seen a c c o r d i n g to the grain size: (i) the twin configuration and (ii) the twin volume fraction. (i) For grade S G , the grains mainly exhibit one twinning system (grains A , Β or C for example, fig. 4). F o r grade L G , m o s t of the grains exhibit 2 twinning systems with t w o different configurations. In grain D , the t w o systems w e r e sequentially activated. A s twin boundaries are
84
·
Materials Processing and Texture
G r a i n S i z e Effect o n T e n s i l e B e h a v i o r of a F e M n C T W I P S t e e l
strong obstacles, the second twin system is blocked by the first o n e , leading to a 'ladder-like' pattern. This configuration occurs in a large n u m b e r o f grains, a s also reported by A l l a i n on the basis o f optical m i c r o g r a p h s . It m a y p r o m o t e the high strain hardening rate in case of grade L G at high deformation level (see f i g . l ) . (ii) T h e v o l u m e fraction o f twins increases with the grain size. W e estimate that the twin v o l u m e fraction at e =0.52 is about 7 % for grade S G and 1 2 % for grade L G (this estimation w a s obtained by a c o m b i n e d analysis of orientation data obtained by E B S D and X R D ) . In fact, for grade L G , the n u m b e r o f t w i n b u n d l e s within t h e grains is h i g h e r a n d the t w i n s are thicker. These two effects lead to a higher twin v o l u m e fraction. The twin orientations o f grains close to the < 1 1 1 > / / R D fiber (in blue) belong to the < 1 0 0 > / / R D fiber (in red). In case o f grade L G , the v o l u m e fraction o f t w i n s s e e m s sufficient to influence the bulk texture. In fact, the v o l u m e fraction o f twin orientations are higher at ε,=0.52 (see figure 3) and consequently the texture less sharp. A d d i t i o n a l analyses h a v e b e e n performed to characterize the onset of twinning during the deformation o f these t w o grades. A l l a i n h a s s h o w n t h a n k s to optical micrographs that twinning started at St=0.10 in grade L G . W e performed T E M analyzes o n grade S G and showed that twinning occurred at 8 = 0 . 0 2 . This confirms that t h e deformation onset for twinning corresponds effectively to the increase o f the strain h a r d e n i n g rate (see figure 1). Moreover, the corresponding stress levels are close (about 5 3 0 M P a ) . I n the literature, the d e p e n d e n c e o f the twin initiation stress o n the grain size h a s not been adequately q u a n t i f i e d . In t h e studied range o f grain size, the critical stress for twin nucleation seems not grain size dependent. Recently Bouaziz et a l . p r o p o s e d a n e w physical m o d e l , adapted from the previous w o r k of Sinclair , to describe the effect o f grain size on the w o r k h a r d e n i n g m e c h a n i s m s o f T W I P steels. This model also suggests that the deformation onset for twinning decreases with the grain size whereas the corresponding stress values remain constant. 13
t
13
1 4
t
4,6
15
16
SUMMARY T h e tensile behavior o f high m a n g a n e s e austenitic T W I P steels has been investigated in relation t o microstructure and texture evolutions. A r e m a r k a b l e c o m b i n a t i o n o f high strength and high ductility w a s observed, mainly d u e t o a high h a r d e n i n g rate u p to large strains. In fact, the mechanical t w i n s act as obstacles for dislocation gliding ( d y n a m i c Hall Petch effect) ' and promote this high h a r d e n i n g rate. 7
8
T h e tensile deformation p r o d u c e s a sharp austenitic texture that strongly influences the twinning activation. In fact, a strong < 1 1 1 > / / R D fiber texture and a w e a k < 1 0 0 > / / R D fiber develops progressively with the deformation, mainly d u e to grain rotations induced by slip. Grains in < 1 1 1 > / / R D orientations activate m o r e t w i n s then those in < 1 0 0 > / / R D orientations, in agreement with S c h m i d Factor calculations. T h u s the texture evolution progressively increases the n u m b e r o f grains favorably oriented for twinning. The corresponding microstructure is rather c o m p l e x with t w i n s organized in stacks, or bundles, of n a n o - t w i n s . T h e grain size (for variation b e t w e e n 2 . 6 μ π ι to 15μπι) also strongly influences t h e twin activity a n d consequently the texture evolution and the resulting mechanical behavior. T h e following r e m a r k s c a n be m a d e c o n c e r n i n g the grain size effect: (i) T h e m a c r o s c o p i c stress onset for t w i n n i n g s e e m s grain size independent (in the studied range of grain size) and is about 5 3 0 M P a .
Materials Processing a n d Texture
·
85
G r a i n S i z e Effect o n T e n s i l e B e h a v i o r of a F e M n C T W I P S t e e l
(ii) According to the Hall-Petch relation, this stress level is observed for lower strains in case of small grain-sized grades, leading to a higher strain hardening rate at low strain levels (due to twin activation). (iii) For large strain levels, the v o l u m e fraction of twins is higher in large grain-sized grades: the twins are thicker - the n u m b e r of t w i n bundles within the grains is higher - t w o t w i n systems are present in most of the grains. In fact in small grain-sized grades, grains <1μηι exhibit no t w i n s (strong grain size effect) and g r a i n s > ^ m exhibit mainly one twin system. T h e higher v o l u m e fraction of twins in a ' t w o - s y s t e m ' configuration p r o m o t e s a higher strain h a r d e n i n g level at large strain levels, for large grain-sized grades. ACKNOWLEDGMENT This w o r k is achieved in the Technological Research T e a m ' M i s e en forme des a c i e r s ' . T h e authors thank all the researcher involved in this team, and ArcelorMittal w h o provided the material. T h e austenitic F e M n C T W I P steels are the result of a j o i n t venture w o r k b e t w e e n ArcelorMittal and T h y s s e n K r u p p Stahl. REFERENCES I. Karaman, H. Sehitoglu, A. J. Beaudoin, Y. I. C h u m l y a k o v , H.J. Maier, C. N . T o m é , Acta Materialia, 4 8 , p . 2 0 3 1 - 2 0 4 7 (2000) L. R é m y , Kinetics, Acta Metallurgica, 2 6 , p.443-451 (1978) S. Asgari, E. El-Dabaf, S.R. Kalidindi, R.D. Doherty, Metallurgical and Materials Transactions A, 2 8 A , p . 1 7 8 1 - 1 7 9 5 (1997) E. El-Danaf, S.R. Kalidindi, R . D . Doherty, Metallurgical and Materials Transactions A, 3 0 A , p.1223-1233 (1999) P. Y a n g , Q. X i e , L. M e n g , H. D i n g b , Ζ . T a n g b , Scripta Materialia, 5 5 , p . 6 2 9 - 6 3 1 (2006) C. Scott, N . Guelton, S. Allain, M . Farral, Proceeding of the MS&TO5 Conference (2005) H. J. B u n g e , Texture Analysis in Material Sciences, ed. D G M ( 1 9 8 2 ) S. Mahajan, G. Y. Chin, Acta Metallurgica, V o l u m e 2 1 , Issue 2, p . 1 7 3 - 1 7 9 (1973) L. R e m y , Acta Metallurgica, V o l u m e 2 5 , Issue 6, p . 7 1 1 - 7 1 4 (1977) B . Petit, N . G e y , M . Cherkaoui, B . Bolle, M . Humbert, International Journal of Plasticity, 2 3 2, p . 323-341 (2007) D . Barbier, Ν . G e y , Β . Bolle, Μ . H u m b e r t , S. Allain, Ο. Bouaziz, Proceeding Matériaux 2006 (2006) S. V e r c a m m e n , B . Blanpain, B . C . D e C o o m a n , P. Wollants, Acta Materialia, 52, 7, p . 2 0 0 5 2012(2004) S. Allain, J. P. Chateau, Ο. Bouaziz, M . Legros, X . Garât, International Conference on TRIPAided High Strength Ferrous Alloys, p . 7 5 - 7 7 (2002) D . Barbier, Ν . G e y , M . H u m b e r t , N . B o z z o l o , S. Allain, to be published Ο . Bouaziz, S. Allain, C. Scott, Scripta Materialia, article in press (2007) C . W Sinclair, W . L . Poole, Y. Bréchet, Scripta Materialia, 5 5 , p . 739-742 (2006) 1
2
3
4
5
6
7
8
9
1 0
11
1 2
1 3
1 4
1 5
1 6
86
Materials Processing and Texture
COMBINATION OF T W O X-RAY DIFFRACTION SETTINGS FOR DETERMINING THE TEXTURE OF SMALL SHEARED SAMPLES 1
2
1
1
D. Barbier , Β . B o l l e , J. J. F u n d e n b e r g e r , C. L a m e l l e , A . T i d u
2
1
Laboratoire d ' E t u d e d e s Textures et Applications aux Matériaux ( L E T A M ) , C N R S U M R 7078, Université Paul Verlaine M e t z , Ile d u Saulcy B P 8 0 7 9 4 , 5 7 0 1 2 M e t z C e d e x , France Laboratoire d ' E t u d e des Textures et Applications a u x Matériaux ( L E T A M ) , C N R S U M R 7 0 7 8 , Ecole Nationale d ' I n g é n i e u r s d e M e t z , Ile du Saulcy, 57045 M e t z C e d e x , France 2
ABSTRACT W e developed a method to perform accurate texture determination on small size sample. Limitations o c c u r w h e n samples having a small g a u g e area are investigated by classical X R a y diffraction. A s t h e irradiated surface increases w i t h the tilt a n g l e , it b e c o m e s larger t h a n the gauge area. T h e m e a s u r e d pole figures are then reduced to low values of the tilt angle and do not allow an accurate texture analyze. T o enlarge the m e a s u r e m e n t range of pole figures and perform m e a s u r e m e n t s o n small samples, w e have used t w o c o m p l e m e n t a r y m e t h o d s . T h e first one is the Bragg-Brentano geometry w h i c h is used for small tilt angles. In the second one, used for larger tilt angles, the size of the irradiated surface by the incident X R a y b e a m remains lower than the sample surface t h a n k s to a specific diffraction setting developed for thin film measurement. The diffraction experiment is conducted so that the incidence angle b e t w e e n the incident beam and the sample surface is kept constant. Consequently, the irradiated area is constant during the measurement. T h e choice of the convenient incidence angle, w h i c h permits to extend the pole figure m e a s u r e m e n t s for high tilt angles is then an important parameter. T h e pole figures are reconstructed from the intensities obtained by these different m e a s u r e m e n t s . In this contribution, we present in details these c o m b i n e d techniques applied to the texture m e a s u r e m e n t of a high strength F e M n C austenitic T W I P steel during simple and reverse shear testing. INTRODUCTION In the field of the d e v e l o p m e n t of high strength alloy for automotive industry, the quantitative texture analyze during the deformation of the material is essential for the understanding o f it behavior. A s the metal sheet undergoes c o m p l e x deformation during the forming process, several devices w e r e developed in the recent year to reproduce these conditions by the c o m b i n a t i o n s of several mechanical tests, like the reverse shear or the strain path change (plane-strain / simple s h e a r ) . M o s t of the time, the g a u g e area of these samples is limited to a few millimeters width ( 3 . 5 m m in o u r study) w i t h a sheet thickness of nearly 1 m m . T h i s small size conducts to the decrease of the accuracy of the texture analyze. In fact, texture m e a s u r e m e n t limitations occur w h e n small size samples (ie samples having a small gauge area) are investigated by classical X - R a y diffraction. The dimension of the irradiated surface increases with the goniometer tilt angle (ψ) and can b e c o m e larger than the gauge surface. T h e analyzed P F s (PF) are then reduced to low values of the tilt angle ψ and do not allow an accurate texture analyze. In o u r study for example, the PF m e a s u r e m e n t is limited to ψ = 6 0 ° for t h e {111} P F . M o r e o v e r deformation b y shear c o n d u c t s to P F s h a v i n g intensity at their border, ie at χ > 5 0 ° . This information will not b e accessible by using classical X-Ray diffraction geometry. Therefore, to accurately study the texture evolutions of small samples submitted to shear tests, the use of adapted X - R a y techniques is necessary. 1,2
3
87
C o m b i n a t i o n of T w o X - R a y Diffraction S e t t i n g s for D e t e r m i n i n g T e x t u r e
EXPERIMENTAL TECHNIQUE A classical method to m e a s u r e P F s with X-ray diffraction is the Schultz o n e , with the B r a g g - B r e n t a n o ( B B ) geometry. O n e of the c o n s e q u e n c e s of this geometry is that the size of the irradiated surface increases with the g o n i o m e t e r tilt angle ψ. W i t h a cylindrical collimator, the irradiated surface has an ellipsoidal shape, characterized by the length L o f it major a x i s (figure 1)· 4
Figure 1. Irradiated surface geometry (collimator diameter d; greatest size o f the irradiated surface L ; real incidence angle i). The greatest size L of t h e ellipse d e p e n d s o n the m e a s u r e m e n t conditions, ie the real incidence angle i itself related to 9 i and ψ by the following relation (1 ): nc
sin i = sin 0 j
nc
cos ψ
(1)
During a classical P F m e a s u r e m e n t (9m constant), w h e n the g o n i o m e t e r tilt angle ψ increases, t h e real incidence angle i decreases, resulting in a n a u g m e n t a t i o n of the length L . O n the figure 2, the evolution of L is plotted with i for a collimator diameter of 0 . 8 m m . For l o w values o f i, ie high values ψ, L b e c o m e s very large. For a {111} P F , L b e c o m e s larger t h a n the smallest d i m e n s i o n of t h e sample g a u g e area (L>3.5 m m ) from i = 1 3 ° , ie ψ=6Ό°. C
F i g u r e 2. Evolution of the size L of the irradiated surface with the real incidence angle i, for the {111} PF with a collimator diameter of 0.8 m m .
In the present case, w h e n L is larger than the sheared zone (L>3.5 m m ) , the b e a m will overlap both the clamping and the g a u g e areas. T h e analyzed surface will not be representative from the deformed one. M o r e o v e r the homogeneity of the deformation is limited to a restricted part of t h e gauge area. T o increase the analyzed surface o v e r 3.5 m m in width, a n d g o through the problem link to the B r a g g - B r e n t a n o geometry, the sample is classically cut into t w o pieces
88
·
Materials Processing and Texture
C o m b i n a t i o n of T w o X - R a y Diffraction S e t t i n g s for D e t e r m i n i n g T e x t u r e
w h i c h are then p l a c e d side by side. B u t in these conditions, the cutting of the s a m p l e should be d o n e very precisely and the deformation will not be h o m o g e n o u s in the m e a s u r e d area. S o , to perform fine texture m e a s u r e m e n t s o n such s a m p l e s , the B B geometry is inappropriate because the m e a s u r e m e n t s are limited to l o w values of the tilt angle. T h e n , to collect data for the high v a l u e s of the tilt angle, t h e greatest size L o f the irradiated surface has then to b e limited. T o k e e p the length L constant, t h e real incidence angle i should not increase a n y m o r e . This c a n b e achieved t h a n k s t o a technique d e v e l o p e d by S z p u n a r ' and P l a y e r for the texture m e a s u r e m e n t o n thin films. T h e w o r k i n g o u t o f this m e t h o d requires to specifically drive t w o angles o f the g o n i o m e t e r (9i„ , ψ ) to k e e p the real incidence angle i constant. T h e goniometer angle a n d the g o n i o m e t e r tilt angle ψ should satisfy the relation (1) with the condition sin(i)=constant. T h e a n g l e s are given on the figure 4 . 5
6
7
c
Figure 4. Diffraction g e o m e t r y a n d definition o f the angles. T h e a n g l e s are defined as followed: 9i„ : g o n i o m e t e r incident b e a m a n g l e ψ: g o n i o m e t e r tilt angle φ : g o n i o m e t e r a z i m u t h angle c
9 : B r a g g angle i: real incidence angle χ: P F tilt a n g l e Φ : P F azimuth angle D
ΰ)=θ -θ ίηο
Β
B u t this g e o m e t r y h a s several c o n s e q u e n c e s : (i) T h e limitation of the P F c o v e r a g e to the P F tilt angle defined by the equations (2) and (3): Z „=|0B-i|
(2)
M I
X
m
a
x
( 1 sin i = a c o s - — 1 2 sin 9 B
J
... (3)
Materials Processing and Texture
89
C o m b i n a t i o n of T w o X - R a y Diffraction S e t t i n g s for D e t e r m i n i n g T e x t u r e
(ii) T h e correction of the intensities from the diffracted v o l u m e . The intensities are corrected considering t h e evolution of the diffracting according to the relation (4) [8]: Γ Γ = 1 = 1 —tan ω (4) tan9
volume
B
With: μ: linear absorption coefficient IBB: diffracted intensities with the Bragg-Brentano geometry (iii) T h e localization of the measured intensities in the PF reference frame. The tilt and azimuth angles from the PF are expressed according to (5) and ( 6 ) : 8
cos χ = cos lycos ω
(5)
.*. A ·.!.·• sin 05 .,. Φ = φ + Δφ with smA(f = (6) sin χ F r o m these relations, the m e a s u r e d intensities are replaced in the PF reference frame. The choice of the convenient incidence angle, w h i c h permits to extend the PF m e a s u r e m e n t s for h i g h tilt angles, is t h e n an important parameter and will limit the size of the irradiated surface. B y using the constant incidence g e o m e t r y and with the appropriate intensities corrections, w e are able n o w to collect and analyze data for the h i g h values of the tilt angle. A s this geometry limits the PF c o v e r a g e to a range of high tilt angles, the resulting blind area at the PF center will be m e a s u r e d by using the classical B B geometry. The global P F is then obtained by c o m b i n i n g the one m e a s u r e d with the B B geometry for the low tilt angles and the one obtained with t h e constant incidence technique for the high tilt angles. W e applied this technique to an austenitic F e M n C a l l o y w h i c h was recently developed for n e w automotive applications. In this paper, w e will only present the results obtained o n the {111} P F s . 9
VALIDATION OF THE TECHNIQUE T o validate the m e t h o d , w e first applied this technique to a large size sample deformed by uni-axial tension. A s the gauge area of the sample is 2 0 m m in width, both techniques can be applied. Different m e a s u r e m e n t s using different values of the real incident angle ( 1 7 ° , 13° and 9°) are c o m p a r e d to the reference one performed with the B B geometry. The figure 5 presents the correction of the intensities localization for the {111} PF measured with a n incident angle of 17°.
90
·
Materials Processing and Texture
C o m b i n a t i o n of T w o X - R a y Diffraction S e t t i n g s for D e t e r m i n i n g T e x t u r e
Figure 5. Correction o n the intensities localization o n the {111} P F for an incident angle of 17°. Without the correction (a) and after t h e correction (b).Levels: - 9845.43-
T h e P F s are then corrected from the variation of the diffracted volume and from the background noise. T h e final results are presented on the figure 6. T h e P F s obtained with the three constant incidence angles reproduce with a very good accuracy the o n e measured with the B B geometry. T h e intensities and their locations are well reproduced from one geometry to another.
b)
Figure 6. C o m p a r i s o n of the {111 j P F s obtained with the B B geometry and the constant incidence one using an angle of 17°, 13°, and 9 ° . L e v e l s : -7973 For the P F m e a s u r e d with a constant incident angle, the blind area in the center of the P F s increases with the decrease o f the real incidence angle T, reaching a range of 25° for i=9°.
Materials Processing and Texture
91
C o m b i n a t i o n of T w o X - R a y Diffraction S e t t i n g s for D e t e r m i n i n g T e x t u r e
EXPLOITATION This technique is n o w applied to a sheared sample on the s a m e material which w a s deformed to a shear strain o f γ = 0 . 6 , with a g a u g e area of 3.5 m m . T h e first step consists in the c h o i c e of the convenient real incidence angle. In the figure 1, w e reported in the detailed graph t h e upper limit of the g a u g e area of the sheared sample. T o k e e p the length L o f the irradiated surface lower than this value, ie 3.5 m m , the limit value of the real incidence angle is i = 1 3 ° . U n d e r this value, L will be larger than 3.5 m m . W e chose i=13° to avoid any overlap with the c l a m p i n g areas. The m e a s u r e m e n t s w e r e then performed o n a S I E M E N S D 5 0 0 0 goniometer u s i n g a using a Cr K„ radiation with a cylindrical collimator with a diameter of 0.8 m m . In these conditions, three P F s are m e a s u r e d : {200} {220} and {111} (only the results o n the {111} PF are presented). T h e details for the P F tilt angle range for each geometry are reported in the table 1. Table 1 : measurement conditions for each geometry a n d P F s . 111 A
200 A
220 A
χ for the BB geometry
0—60
0—65
0—80
χ for a real incident angle of 13°
30->77.5
35—77.5
60—82.5
For each P F s , w e check the positions and the values of the intensities in the overlapping zone b e t w e e n the t w o geometries, to analyze the coherence of the t w o m e a s u r e m e n t s . For e x a m p l e , this z o n e goes from χ = 3 0 ° to χ = 6 0 ° for the {111} P F . T h e figure 7a correspond to the measured {111} P F with the B B g e o m e t r y , the 7b with the constant incidence technique using a real incidence angle of 13°. F r o m these t w o m e a s u r e m e n t s , the global P F is then reconstructed (figure 7c).
Figure 7. M e a s u r e d {111} P F (a) w i t h the B B geometry (b) with the constant incident angle m e t h o d (c) global P F obtained with (a) and (b). L e v e l s : W i t h the classical geometry (ie the B B ) , t h e {111} P F (figure 7a) is limited to χ = 6 0 ° . By using the combination of the B B and the constant incidence geometry, it is possible to perform m e a s u r e m e n t u p to χ=77.5° (figure 7c).
92
·
Materials Processing and Texture
C o m b i n a t i o n of T w o X - R a y Diffraction S e t t i n g s for D e t e r m i n i n g T e x t u r e
SUMMARY A m e t h o d to m e a s u r e the crystallographic texture of sheared sample with a gauge area under 4 m m has been developed. This method c o m b i n e s to different m e a s u r e m e n t geometries to increase the P F coverage. The technique w a s first validated o n large size sample and then used to characterize the texture evolution o f a T W I P steel during shearing. By using the resulting P F s , with a larger coverage, the determination of the orientation distribution function is improved, w h i c h permits a better understanding of the deformations m e c h a n i s m s of this steel. T h i s m e t h o d can be applied in every cases w e r e the size of the analyzed area is restricted. But the m e a s u r e d area has to be statistically representative of the material. REFERENCES E.F. R a u c h and J.-H Schmitt, Mater. Sei. Eng. A, v o l . 1 1 3 , p p . 4 4 1 - 4 4 8 (1989) A . B . L o p e s , E.F. R a u c h and J.J. Gracio, Λεία Mater., vol.47 (3), p p . 859-866 (1999) G.R. C a n o v a , U . F . K o c k s and J.J. C a n o v a , Acta Metall., vol.32 (2), p p . 2 1 1 - 2 2 6 (1984) L.G. Schulz, J. Appl. Phys., vol.20, p p . 1030-1033 (1949) J A . Szpunar, S. A h l r o o s and P. Tavernier, J. Mat. Sei., v o l . 2 8 , pp. 2 3 6 6 - 2 3 7 6 (1993) J.A. Szpunar and P. Blandford, Mat. Sei. Forum, vol. 157-162, p p . 2 0 7 - 2 1 2 (1994) M . A. Player, G. V. Marr, E. Gu, H. Savaloni, N . Ö n c a n and I. H. M u n r o , J. Appl. Cryst., vol.25, p p . 7 7 0 - 7 7 7 ( 1 9 9 2 ) J.J. H e i z m a n n and C. L a m e l l e , J. Appl. Cryst., vol.19, pp 4 6 7 - 4 7 2 (1986) C. Scott, N . Guelton, S. Allain, M . Farral, T h e development of a n e w F e - M n - C austenitic T W I P steel for automotive applications, Proceeding of the MS&TO5 Conference (2005) 1
2
3
4
5
6
7
8
9
Materials Processing and Texture
·
93
This Page Intentionally Left Blank
TEXTURE OF DIFFRACTION
NON-ORIENTED
ROLLED
STEELS
BY
X-RAY
AND
ELECTRON
Martin Cernik, MiloS P r e d m e r s k y , Andrej Leäko, Erik Hilinski U. S. Steel Kosice, s.r.o., Vstupny A r e a l U. S. Steel, U S S E Research, Metallography Department Koäice, Slovak Republic ABSTRACT A c o m p a r i s o n of different texture m e a s u r e m e n t m e t h o d s including X - R a y Diffraction ( X R D ) and Electron Backscattering Diffraction ( E B S D ) , and orientation distribution function (ODF) calculation algorithms, W I M V , A D C , and S H , w a s m a d e o n t w o non-oriented electrical steels. T h e first steel w a s l o w silicon steel that w a s subjected to t w o hot rolling practices. Texture analysis w a s accomplished via O D F s and (100)/(111) crystal plane v o l u m e fraction ratios, also termed the magnetic texture quality ratio, determined from the W I M V , A D C , and E B S D analyses. It w a s determined that the all m e a s u r e m e n t techniques and O D F calculation algorithms are equally good and reliable and that the (100)/(111) v o l u m e fraction ratio is a good measure for characterizing the quality of electrical steel textures. A s a result of these findings only the E B S D technique w a s used to evaluate the effect of Sb on texture development in a second high silicon steel annealed in a laboratory continuous annealing simulator. In that study the ( 1 0 0 ) / ( U 1 ) crystal plane v o l u m e fraction analysis s h o w e d that S b improved magnetic properties w h e n the final grain growth annealing temperature w a s 1000°C or less. INTRODUCTION T h e r e are three major experimental techniques used to determine the degree of preferred orientation in materials: n e u t r o n , X - R a y Diffraction ( X R D ) and Electron Backscatter H o w e v e r because of accessability, most researchers opt to use either Diffraction ( E B S D ) . X R D or E B S D m e a s u r e m e n t techniques. In the X R D technique, a texture goniometer is used to measure pole figures w h i c h are in turn used to determine the orientation distribution function ( O D F ) . In the E B S D method, the O D F s are calculated directly from the m e a s u r e d inverse pole figures (IPFs). T h e O D F qualitatively and quantitatively describes the textural state of material in the excited v o l u m e of material. T h e O D F calculation can b e performed via several mathematical procedures, w h i c h are included in commercially available texture analysis software. M o r e specifically, those explored in this paper are the W I M V a l g o r i t h m in p o p L A , the A D C m e t h o d in L a b o T e x , and B u n g e ' s S H m e t h o d in the T S L O I M software package. E a c h of the a b o v e m e n t i o n e d texture m e a s u r e m e n t m e t h o d s and analysis software has advantages and disadvantages, w h i c h depend to a great extent o n the subject material and its texture. F o r rolled steel sheets both X R D and E B S D texture m e a s u r e m e n t m e t h o d s have been c o m m o n l y e m p l o y e d . X R D ' s advantages are in easy sample preparation, m o r e penetrable radiation, a higher sampling v o l u m e , and faster analyses. While E B S D sample preparation is more d e m a n d i n g , the acquired data links the microtexture w i t h the microstructure providing more useful information to the researcher. T h e m a g n e t i c properties o f electrical steels are well k n o w n to depend upon the preferred orientation of the g r a i n s and thus serve as a good material to c o m p a r e the X R D and E B S D based texture m e a s u r e m e n t m e t h o d s . W h e r e a s most authors use one texture analysis method, this paper presents a comparison of texture analyses obtained from W I M V , SH, and A D C generated O D F s o n a commercially produced l o w silicon steel subjected to t w o hot rolling 1
2 - 3
4,5
6,7
8
9
10
95
T e x t u r e of N o n - O r i e n t e d R o l l e d S t e e l s b y X - R a y a n d E l e c t r o n Diffraction
practices. The texture analysis m e t h o d s developed are then used to explore the effect of S b on final texture in experimental h i g h silicon electrical steel.
MATERIAL AND EXPERIMENTS The evolution of texture in t w o electrical steels w a s investigated to c o m p a r e textures measured using X R D and E B S D techniques. In the first steel, a commercially p r o d u c e d l o w Si M 8 0 0 - 6 5 H P grade, modifications to the hot rolling practice w e r e explored. T h e final texture resulting from the modified austenite/ferrite finish hot rolling practice (sample 1) w a s compared to that obtained from the standard full ferrite finish rolling (sample 2). Their compositions and resultant magnetic properties are listed in Table I. The texture analysis techniques developed to characterize the M 8 0 0 - 6 5 H P w e r e then used to determine the influence of antimony (Sb) o n texture evolution in an experimental degassed M 4 0 0 - 5 0 A grade, subjected to final annealing in a laboratory continuous annealing simulator. Commercially cold rolled sheet, with (sample B ) and without (sample A ) S b , were secured and laboratory decarburized at either 830 or 9 0 0 ° C followed by grain growth annealing at temperatures ranging from 9 0 0 to 1060°C, as indicated in Figure 1. T h e respective chemical compositions and magnetic properties are s h o w n in Table II.
Table I. C o m p o s i t i o n (wt. % ) and magnetic properties of the M 8 0 0 - 6 5 H P l o w silicon steel Sample
C[%]
Mn [%]
Si [%]
1 2
0.0047 0.0089
0.198 0.040
0.636 0.610
96
·
0.015 0.120
S[%]
Al [%]
Ν [%]
Jsooo [Τ]
Pis [W/kg]
ASTM Grain Size No.
0.0058 0.0057
0.0023 0.0030
0.0025 0.0028
1.778 1.721
6.475 7.630
7.5 8.5
Materials Processing and Texture
T e x t u r e of N o n - O r i e n t e d R o l l e d S t e e l s b y X - R a y a n d E l e c t r o n Diffraction
Table II. Composition (wt. % ) and m a g n e t i c properties of the M 4 0 0 - 5 0 A high silicon steel Sample
C[%]
Mn [%]
Si [%]
P[%]
S[%]
Al [%]
Sb [%)
A Β
0.007 0.007
0.25 0.25
2.4 2.4
0.005 0.007
0.0043 0.0041
0.40 0.40
0.06
J 5000
[T]
1.670 1.690
Pl.5 [W/kg] 3.287 3.174
ASTM Grain Size No. 5.0 4.5
F o r the X R D texture studies, flat s a m p l e s 30 χ 30 m m in size were prepared by grinding followed by chemical etching in a 7 % H F in H 2 O 2 solution to r e m o v e residual deformation left from grinding. A Seifert X R D 3 0 0 3 diffractometer equipped with a texture goniometer and using M o Κ α radiation w a s used to m e a s u r e the (110), (200), (112) and (103) incomplete pole figures in back reflection m o d e . O D F s w e r e calculated via the p o p L A and L a b o T e x programs, employing the W I M V and A D C iteration m e t h o d s , respectively. In addition, the A D C algorithm w a s also used to p r o v i d e the v o l u m e fraction of selected texture c o m p o n e n t s . S a m p l e s for texture analysis via E B S D w e r e prepared b y grinding, polishing with 1 μηι diamond paste, followed b y further vibration polishing in a 50 n m colloidal solution. Textures were m e a s u r e d in a FEI Q u a n t a 4 0 0 tungsten filament S E M operating at 25 k V with a C C D E B S D c a m e r a . IPF m a p s w e r e m e a s u r e d and the O D F s calculated directly from the m a p s using B u n g e ' s S H functions provided in the T S L O I M software. A s a matter of note, the sample preparation p r o c e d u r e s for E B S D w e r e m o r e laborious and to improve the statistics necessary to reproduce the macrotexture results from the X R D up to 10 individual fields required measurement o n one sample. RESULTS AND DISCUSION T h e description of electrical steel textures in simplified form c a n b e accomplished by considering the η , α and γ-fibers, w h e r e the η-fiber is defined by {1.00}//ND, α by < 1 1 0 > / / R D , and γ by {111 }//ND. In addition the v o l u m e fractions o f the m a i n texture c o m p o n e n t s are also useful for characterization. T h e v o l u m e fraction of material, w h e r e the crystallites are oriented in the < h k l > crystallographic direction parallel to the sheet normal direction ( N D ) is calculated using the O D F f(g):
V
=
i
fi8) 8
()
T h e texture of samples 1 and 2, that c o m p a r e hot rolling practices are characterized via the O D F s . T h e ψ2 = 45° O D F sections calculated from X R D m e a s u r e d pole figures using p o p L A and L a b o T e x are s h o w n in Figures 2 a through 3b. L i k e w i s e the s a m e O D F sections for samples 1 and 2 calculated from the E B S D IPF m a p s are s h o w n in Figures 4 a and 4 b . T h e r e is good consistency b e t w e e n the three m e a s u r e m e n t sets yet there are some subtle differences. T h e X R D m e a s u r e m e n t c o m b i n e d with the W I M V and A D C algorithms tend to accentuate intensity and reveal m o r e detail along the η-fiber. T h e E B S D measurements in conjunction with S H algorithm generally soften the overall intensity while showing m o r e Goss. E x a m i n i n g the three m e a s u r e m e n t sets, it can b e concluded that sample 1 exhibits a higher portion of G o s s ( 0 1 1 ) < 1 0 0 > and rotated cube texture c o m p o n e n t s w h e r e a s sample 2 has a texture defined by a stronger γ-fïber with a ( 1 1 1 ) < 1 - 2 1 > texture c o m p o n e n t . T h e magnetic quality of the texture can b e characterized by the simple v o l u m e fraction ratio (100)/(111) of the respective
Materials Processing and Texture
97
T e x t u r e of N o n - O r i e n t e d R o l l e d S t e e l s b y X - R a y a n d E l e c t r o n Diffraction
98
Figure 2a. O D F of s a m p l e 1 calculated from p o p L A ( W I M V ) .
Figure 2 b . O D F o f s a m p l e 2 calculated from p o p L A ( W I M V ) .
Figure 3a. O D F of sample 1 calculated from L a b o T e x ( A D C ) .
Figure 3 b . O D F of sample 2 calculated from L a b o T e x ( A D C ) .
Figure 4a. O D F of sample 1 calculated from T S L O I M ( S H ) .
Figure 4 b . O D F of s a m p l e 2 calculated from T S L O I M ( S H ) .
Figure 5a. IPF m a p of s a m p l e 1 via E B S D .
Figure 5b. IPF m a p of sample 2 via E B S D .
·
Materials Processing and Texture
T e x t u r e of N o n - O r i e n t e d R o l l e d S t e e l s b y X - R a y a n d E l e c t r o n Diffraction
texture c o m p o n e n t s . F o r s a m p l e 1 this ratio is 0.40 w h e r e a s for sample 2 it is 0.32 for the WTMV O D F s . T h e A D C m e t h o d yielded respective ratios of 0.75 and 0.24. Simplified I P F m a p s for the three basic crystallographic directions in samples 1 and 2 are given in Figures 5a and 5 b . A s w a s d o n e with t h e O D F s , the m a g n e t i c texture quality ratio w a s determined from the IPFs. T h e values are 0.45 and 0 . 2 3 . respectively which are in good agreement with those obtained from the W I M V O D F s . Therefore, t h e (100)/(111) volume fraction ratio, serves as a simple quality m e a s u r e for electrical steel textures. It is obvious that magnetic quality can b e i m p r o v e d b y increasing the a m o u n t of (100) o r reducing (111) textures. A n analysis o f the O D F s for s a m p l e s 1 and 2 s h o w s that the superior texture obtained from the modified hot rolling practice given sample 1 w a s due to a w e a k e n i n g of the γ-fiber, i.e. a lower volume fraction of (111) planes. T h u s the modified austenite/ferrite finish hot rolling practice delivers m o r e favorable m a g n e t i c textures, as taught b y K a w a m a t a et a l . " and Fischer and S c h n i e d e r and superior m a g n e t i c properties as presented in Table I. 12
B e c a u s e of t h e good a g r e e m e n t obtained between the X R D and E B S D texture evaluations of the M 8 0 0 - 6 5 H P h o t rolled with the standard and modified practices, only the E B S D technique w a s used to characterize the influence of S b on texture evolution during grain growth annealing of a high Si M 4 0 0 - 5 0 A non-oriented electrical steel. Figures 6a and 6 b c o m p a r e the I P F m a p s of samples A and B . It m u s t b e noted that because of the grain size (about A S T M N o . 4.5 - 5) 10 measurement fields, such as s h o w n in Figures 6, w e r e used to improve the statistical accuracy of the O D F s . Stitching together the m e a s u r e m e n t fields yielded a total m e a s u r e m e n t area of about 2.5 m m , w h i c h is larger t h a n the m e a s u r e d area in the X R D . For very large grains it is necessary to increase the n u m b e r o f I P F m a p s and the step size to maintain the same measurement time for a sample. 2
Figure 6a. E x a m p l e field of an I P F m a p o f sample A used to m e a s u r e the O D F .
Figure 6b. E x a m p l e field of an I P F m a p of s a m p l e Β used to m e a s u r e the O D F .
As s h o w n in Figures 6 and Table II, the addition o f Sb had n o statistically significant effect on increasing grain size. T h e η and γ-fibers, taken from the O D F s measured o n samples subjected to a decarburization anneal of 8 3 0 ° C followed by grain growth anneals o f 9 0 0 , 9 8 0 , and 1060°C are s h o w n in Figures 7a and 7b for the S b free and in Figures 8a and 8b for the Sb alloyed material. Texture i m p r o v e m e n t d u e to the antimony addition c a n n o t b e seen in the IPF m a p s or the calculated O D F s . In fact, w h i l e the γ-fiber behaviors and intensities are similar, the η-fibers indicate that the Sb free material possesses a slightly superior texture. T h i s is contrary
Materials Processing and Texture
·
99
T e x t u r e of N o n - O r i e n t e d R o l l e d S t e e l s b y X - R a y a n d E l e c t r o n Diffraction
Figure 9a. Effect grain g r o w t h annealing temperature o n the m a g n e t i c texture quality ratio on material decarburized at 8 3 0 ° C .
100
Materials Processing a n d Texture
Figure 9 b . Effect grain g r o w t h annealing temperature on the magnetic texture quality ratio o n material decarburized 9 0 0 ° C .
T e x t u r e of N o n - O r i e n t e d R o l l e d S t e e l s b y X - R a y a n d E l e c t r o n Diffraction
13
to the expected behavior that is well k n o w n and patented by I r i e , and noted by m a n y others and S n lead to texture improvement in non-oriented electrical steels. where Sb Therefore to explore this incongruity m o r e closely, a magnetic texture quality ratio analysis w a s conducted. T h e (100)/(111) v o l u m e fraction ratios as determined from the IPF m a p s for all samples are s h o w n in Figures 9a and 9b for the 830 and 9 0 0 ° C decarburization anneals. Here, t h e effect of Sb on modifying the final texture is m o r e clear. For each decarburization temperature the Sb alloyed material h a s a higher (100)/(111) v o l u m e fraction. H o w e v e r as the grain g r o w t h annealing temperature increases, the impact of S b lessens. This is probably the result of grain growth selection advantages in the S b material being eroded by the grain coarsening p r o c e s s . Figure 9a s h o w s that if using a lower decarburizing temperature, the grain g r o w t h annealing temperatures b e y o n d 1000°C negate the benefit S b whereas with the higher decarburization temperature the benefit of Sb can be extended to grain growth temperatures up to 1060°C. Finally it is w o r t h noting that the i m p r o v e m e n t in magnetic texture led to only slight i m p r o v e m e n t s in m a g n e t i c properties as s h o w n in Table II. That the improvement w a s not m o r e substantial is probably the result of this M 4 0 0 - 5 0 A not being hot b a n d annealed as Shimanka has noted hot b a n d annealing is needed to reveal the benefits of S b . ' 4 - 1 6
1 7
8
CONCLUSIONS T h e following points can b e concluded from the experimental w o r k and texture analysis method c o m p a r i s o n s e m p l o y e d herein: 1) M e a s u r e m e n t and evaluation of textures via O D F s developed from X R D and E B S D techniques are equally good and reliable. In addition so are the (100)/(111) volume fraction ratios, especially b e t w e e n the E B S D IPF and W I M V O D F m e a s u r e m e n t methods. 2) T h e a d v a n t a g e of the X R D technique is that sample preparation is simpler, standard techniques can b e used to calculate O D F s , and s o m e physical properties such as the coefficient of normal anisotropy can b e measured. T h e disadvantage is that texture analysis of individual grains is complicated. 3) T h e a d v a n t a g e of E B S D texture analyses is that detailed investigation of individual grains are possible yielding preferential grain orientations, grain b o u n d a r y character (twinning), texture inhomogeneities, and analyses of multiphase materials (ferrite, retained austenite). The disadvantage is in m o r e complicated and time c o n s u m i n g sample preparation. 4) T h e texture i m p r o v e m e n t in the lower Si non-oriented steel given the modified hot rolling practice w a s the result of a decrease in the v o l u m e fraction of the magnetically poorest (111) planes in t h e final material. 5) T h e positive effect of Sb additions on increasing magnetically favorable textures w a s confirmed after simulations of various continuous annealing processes, although the benefit of Sb w a s negated with grain g r o w t h annealing temperatures b e y o n d 1000°C. The Sb alloyed material possessed higher m a g n e t i c polarization values and lover core loss values than the material w i t h o u t antimony. DISCLAIMER T h e material in this paper is intended for general information only. A n y use of this material in relation to any specific application should be based o n independent examination and verification of its unrestricted availability for such use, and a determination of suitability for the application by professionally qualified personnel. N o license under any United States Steel
Materials Processing and Texture
•
101
T e x t u r e of N o n - O r i e n t e d R o l l e d S t e e l s b y X - R a y a n d E l e c t r o n Diffraction
Corporation patents or other proprietary interest is implied by the publication of this paper. Those m a k i n g use of or relying u p o n the material a s s u m e all risks and liability arising from such u s e or reliance. REFERENCES Ή . J. B u n g e , Three-dimensional Texture Analysis, Inter. Mat. Rev., 3 2 , 2 5 6 - 2 9 1 , (1987). L . G. Shulz, A Direct M e t h o d of D e t e r m i n i n g Preferred Orientation of a Flat Reflection Sample Using a Geiger C o u n t e r X - R a y Spectrometer, J. of App. Phys., 2 0 , 1030-3, (1949). U . F. K o c k s , C. N . T o m e , H. R. W e n k , Texture and Anisotropv. C a m b r i d g e University Press, 1998. A . J. Schwartz, M . K u m a r , and B . L. A d a m s , Electron Backscatter Diffraction in Materials Science. Kluver A c a d e m i c / Plenum Publishers, N e w York, 2 0 0 0 . V . Randle and O. Engler, Introduction to Texture Analysis. C R C Press, B o c a Raton, 2 0 0 0 . S . Matthies and G. W. Vinel, O n the Reproduction of the Orientation Distribution Function of Textured Samples from R e d u c e d Pole Figures U s i n g the C o n c e p t of Conditional Ghost Correction, Phys. Stat. Sol. B, 112, K l 11-K120 (1982). J . S. Kallend, U. F. K o c k s , A. D . Rollett, H. R. Wenk, Operational Texture Analysis, Mat. Sei. and Eng., Km, 1-11,(1991). K . Pawlik, D e t e r m i n a t i o n of the Orientation Distribution Function from Pole Figures in Arbitrarily Defined Cells, Phys. Stat. Sol. B, 134, 4 7 7 , (1986). B u n g e H. J., Texture A n a l y s i s in Materials Science. M a t h e m a t i c a l M e t h o d s . Butterworths, London, (1982). G . L y u d k o v s k y , P. K. Rastogi, and M . Bala, Nonoriented Electrical Steels, J. of Metals, Jan., pp. 18-26, (1986). R. K a w a m a t a , Y. Arita, S. K a n a o , K. H a n z a w a , T. Kubota, and K. M u r a k a m i , Non-oriented Electrical Steel Sheet Excellent in Permeability and M e t h o d of Producing the S a m e , U. S. Patent N o . 6,425,962 B l . 0 . Fischer and J. Schneider, Influence of Deformation Process o n the I m p r o v e m e n t of N o n oriented Electrical Steel, J. ofMagn. andMagn. Mat, 2 5 4 - 2 5 5 , 3 0 2 - 0 6 , (2003). T . Irie, K. M a t s u m u r a , H. N a k a m u r a , H. Shimanaka, and T. Suzuki, M e t h o d of Producing N o n oriented Silicon Steel Sheets H a v i n g an Excellent Electromagnetic Property, U . S. Patent N o . 4,204,890. 2
3
4
5
6
7
8
9
1 0
, 2
, 3
l4
L y u d k o v s k y , Grigory, and Prabhat K. Rastogi, Effect of A n t i m o n y on Recrystallization Behavior and M a g n e t i c Properties of a Nonoriented Silicon Steel, Met. Trans. A, 1 5 A , February, (1984), p p . 2 5 7 - 2 6 0 . , 5
M . Takashima, T. Obara, and T. Kan, Texture I m p r o v e m e n t in High-Permeability Nonoriented Electrical Steel by A n t i m o n y Addition, J. of Mat. Eng. and Perf, 2, 2 4 9 - 5 4 ( 1 9 9 3 ) . F . V o d o p i v e c , F. Marinsek, F. Gresovnik, D . G n i d o v e c , M . Pracek, and M . Jenko, Effect of A n t i m o n y o n Energy Losses in N o n - o r i e n t e d 1.8 Si, 0.3 Al Electrical Steels, J. of Magn. and Magn. Mat., 9 7 , 2 8 1 - 5 (1991). 1 6
1 7
M . G o d e c , L. Kosec, M . J e n k o , R. Mast, and H. J. G r a b k e , Texture D e v e l o p m e n t in Electrical Steels Influenced by Tin Addition, Metallwissenschaft und Technik, 5 1 , 7 0 2 - 5 , (1997). H . Shimanaka, Y. Ito, T. Irie, K. M a t s u m u r a , H. N a k a m u r a , and Y. S h o n o , N o n - O r i e n t e d SiSteels Useful for Energy Efficient Electrical A p p a r a t u s , Energy Efficient Electrical Steels. A S M TMS, 193-204,(1980). I 8
102
Materials Processing and Texture
EFFECT OF COATING TEXTURE ON POWDERING BEHAVIOR OF INDUSTRIALLY P R O D U C E D G A L V A N N E A L E D C O A T I N G O N IF, I F H S A N D H S Q G R A D E S T E E L S A.Chakraborty, R.K.Ray and D.Bhattacharjee Research and D e v e l o p m e n t Division, T A T A S T E E L Jamshedpur, Jharkhand, India, ΡΓΝ-831007
ABSTRACT A m o n g the Z i n c and Z i n c alloy coated steels used in a u t o m o t i v e bodies, galvannealed coated steel sheets are the m o s t widely accepted due to their superior spot weldability, corrosion resistance, paintability and improved resistance against scratch. In this process the preheated steel sheet is first i m m e r s e d in a liquid zinc a l u m i n u m bath w h i c h is called Galvanizing, and then given a post coating heat treatment. T h e entire process is k n o w n as Galvannealing. T h e heat treatment causes the zinc in the coating to inter-diffuse w i t h the substrate iron to form several iron-zinc intermetallic phases, n a m e l y g a m m a (Γ), g a m m a i ( Γ ι ) , delta (δ) and zeta (ζ) w h i c h are subsequently stacked over the steel substrate. T h e present study deals with the effect of the coating texture o n its formability or p o w d e r i n g behavior with respect to three industrially produced galvannealed steels, an Interstitial Free (IF), an Interstitial Free H i g h Strength (IFHS) and a L o w C a r b o n High Strength Quality ( H S Q ) . T h e coating formability decreases and the p o w d e r i n g tendency increases from IF to I F H S to H S Q . A n attempt h a s b e e n m a d e to correlate this p h e n o m e n o n with the observed increasing degree of discontinuity of the {01.3} fiber found in the galvannealed coating textures from the IF to the I F H S to the H S Q steel. T h e increasing a m o u n t s of iron present in the coated layers of the a b o v e three steels in that order are also supposed to contribute to the observed effects. Keywords:
G a l v a n n e a l e d coating texture, delta phase, p o w d e r i n g resistance
INTRODUCTION A l t h o u g h galvannealing of steel surfaces impart m a n y advantages with respect to weldability, corrosion and paintability, galvannealed layers suffer from an inherent d r a w b a c k in the form o f p o o r formability. H i g h e r iron content in the coating h a s been suggested to be one of the reasons for p o o r e r formability. T h e embrittlement of the coating d e p e n d s not only on the iron content and distribution of different Fe-Zn intermetallic p h a s e s ' ' , but also on the crystallographic texture of the coated layer. T h e coating texture m a y control the friction behavior as well as corrosion resistance of galvanized and galvannealed coated steels " . Again, in galvannealed coating, the substrate texture appears to control the formation and g r o w t h of the different F e - Z n intermetallic p h a s e s . N o t m u c h attention has b e e n paid till date to carry out any systematic investigation o n the nature of the substrate and coating textures in industrially produced galvanized and galvannealed steels. A l t h o u g h s o m e preliminary ideas about the texture and of coated galvanized layer h a s b e e n reported by G a i g n a r d and Esling by V i n c e n t et al . , these are by no m e a n s c o m p l e t e and c o m p r e h e n s i v e . M o s t of the a b o v e researches others are based o n studies of galvanized, electrogalvanized and galvannealed coatings produced in a simulator, w h e r e the coating thickness of Zn or thicknesses of individual Fe-Zn intermetallic phases could be controlled by the researchers. O n the other hand, it is extremely difficult to clearly distinguish the different F e - Z n intermetallic c o m p o u n d s in an industrial product. 7
8
13
14
16
1 7 - 1
103
Effect of C o a t i n g T e x t u r e o n P o w d e r i n g B e h a v i o r of Industrially P r o d u c e d G a l v a n n e a l e d C o a t i n g
T h e present investigation deals with effect of coating texture o n its formability with respect to three industrially p r o d u c e d galvannealed steels, an Interstitial Free (IF), an Interstitial Free High Strength (IFHS) and a L o w Carbon H i g h Strength Quality ( H S Q ) . EXPERIMENTAL PROCEDURE T h e chemical compositions of the steels are s h o w n in Table I. Galvannealing operation w a s carried out in the continuous galvanizing and galvannealing line in T A T A S T E E L , Jamshedpur, India. T h e Z n bath a l u m i n u m content w a s kept at a constant level of 0.135wt.%. T h e different processing parameters during galvannealing are given in Table II. Table I. Substrate Steel Composition ( W t . % ) Sample No. IF IFHS HSQ
C
Mn
Al
Ti
Nb
Si
S
Ρ
0.0022 0.0035 0.0900
0.09 0.45 0.85
0.015 0.023 0.042
0.033 0.039 0.027
0.008 0.015 0.050
0.003 0.095 0.020
0.007 0.010 0.004
0.012 0.031 0.017
Ν (ppm) 28 35 54
Table II. Galvannealing Processing P a r a m e t e r s . Sample No. IF IFHS HSQ
Steel Sheet Thickness ( m m ) 0.7 1.0 1.6
Line Speed (meter/min.) 107 62 58
Galvannealing Power (kW) 786 1148 1296
Skin Pass Elongation (%) 0.50 0.79 1.57
Cross-sectional scanning electron microscopy along with X-ray m a p p i n g w a s carried out in a J E O L 6 4 0 0 Scanning Electron M i c r o s c o p e ( S E M ) operated at 15kV. The quantitative depth profiling ( Q D P ) w a s carried out using a L E C O G D S - 8 5 0 A G l o w D i s c h a r g e Optical Emission Spectroscope ( G D O E S ) . X - R a y Diffraction ( X R D ) studies of the samples w e r e carried out in the usual m a n n e r in a PaNalytical X ' p e r t P R O X-ray diffractometer, using C u K a radiation. T h e X ray generator voltage w a s 4 5 k V and the tube current w a s 4 0 m A . Crystallographic texture of the coating w a s determined using a texture goniometer attached to a PaNalytical X ' p e r t P R O X R D m a c h i n e . T h e resolution of the equipment w a s 0.0001°. For the purpose of m e a s u r i n g pole figures, a thoroughly cleaned sample of the galvannealed steel of the size 2 c m x 2 c m w a s put into the s p e c i m e n holder. In the present study (233), (054), (330), (241) and (249) incomplete pole figures w e r e determined experimentally. T h e O D F s w e r e calculated from the pole figures after correction (for defocusing and absorption) and symmetrization using t h e series expansion method with a limit of l m a x = 2 2 using Labotex software. T h e chemical c o m p o s i t i o n s of the coatings w e r e determined by A S T M standard Gravimetric m e t h o d . T h e overall percentages of Fe, Z n and Al w e r e m e a s u r e d using an Inductive Coupled P l a s m a Spectroscope ( I C P S ) , Spectro Analytical Instrument, G m b H version 2.0e/8/88, G e r m a n y . T h e formability of the galvannealed coatings w e r e m e a s u r e d by D o u b l e Olsen Test method. T h e weight loss w a s calculated before a n d after test. T h i s basically m e a s u r e s the p o w d e r i n g behavior of the coating. 1 9 - 2 1
2 2
104
Materials Processing and Texture
Effect of C o a t i n g T e x t u r e o n P o w d e r i n g Behavior of Industrially P r o d u c e d G a l v a n n e a l e d C o a t i n g
RESULTS Table III represents the percentage of Fe present in the galvannealed coatings over IF, I F H S and H S Q grade substrate steels, obtained by A S T M standard Gravimetric m e t h o d . T h e relevant portion of the X R D pattern of the IF sample is s h o w n in Figure 1. T h e X R D patterns taken from the other t w o samples ( I F H S and H S Q ) are quite similar. T h e cross-sectional scanning electron m i c r o g r a p h s of the coatings are s h o w n in part (a) of the Figures 2, 3 and 4 . The corresponding X-ray m a p s for Fe and Z n are represented in parts (b) and (c) of the s a m e Figures. During s a m p l e preparation for S E M study there w a s a distinct possibility of the edges of the coating falling off the substrate or getting d a m a g e d . T o prevent this, a 2 0 0 μ η ι thin copper strip was used as the supporting material. T h e Q D P - G D O E S study results of the galvannealed coatings are s h o w n in Figures 5(a-c). The percentage thicknesses of the different Fe-Zn intermetallic p h a s e s w e r e determined by superimposing the compositional r a n g e s obtained from standard F e - Z n binary p h a s e diagram for the different phases o n the a b o v e profiles. T h e results are given in a tabular form in Table IV. T h e Φ 2 sections of the complete Orientation Distribution Function ( O D F ) of the galvannealed coatings o n I F , I F H S and H S Q grade steels are presented in Figures 6, 8 and 10 respectively. All the above figures appear rather similar, although there are some subtle differences. This has been brought out clearly in Figures 7, 9 and 11, w h e r e the Φ 2 = 0 ° sections of the O D F s of the three samples have been s h o w n o n a magnified scale. T h e weight losses m e a s u r e d by D o u b l e O l s e n Test m e t h o d are depicted in Table V. Table III. Fe Content of the Galvannealed Coating obtained by W e t C h e m i c a l M e t h o d s . Sample No. IF IFHS HSQ
% F e Content 10.35 11.10 14.61
2 T h « a C)
F i g u r e 1. X R D plot of galvannealed coating o n IF sample.
Materials Processing and Texture
·
105
Effect of C o a t i n g T e x t u r e o n P o w d e r i n g B e h a v i o r of Industrially P r o d u c e d G a l v a n n e a l e d C o a t i n g
Figure 2. Cross-sectional (a) S E M M i c r o g r a p h , (b) X - R a y M a p of Fe and X - R a y M a p o f Zn of Galvannealed Coating over Sample N o . IF.
Figure 3 . Cross-sectional (a) S E M M i c r o g r a p h , (b) X - R a y M a p of Fe and X - R a y M a p of Z n of G a l v a n n e a l e d Coating over Sample N o . I F H S .
Figure 4. Cross-sectional (a) S E M M i c r o g r a p h , ( b ) X - R a y M a p of Fe and X - R a y M a p o f Zn of G a l v a n n e a l e d Coating over Sample N o . H S Q .
106
Materials Processing a n d Texture
Effect of C o a t i n g T e x t u r e o n P o w d e r i n g B e h a v i o r of Industrially P r o d u c e d G a l v a n n e a l e d C o a t i n g
Materials Processing a n d Texture
·
107
Effect of C o a t i n g T e x t u r e o n P o w d e r i n g B e h a v i o r of Industrially P r o d u c e d G a l v a n n e a l e d C o a t i n g
Figure 5. Q D P of Galvannealed Coating obtained by G D O E S o f S a m p l e N o . (a) I F , (b) I F H S and (c) H S Q .
Table IV. Percentage Thickness of different F e - Z n Intermetallic P h a s e s present in Galvannealed Coating obtained from Q D P - G D O E S Profiles. A m o u n t of P h a s e (%)
IF
IFHS
HSQ
Γ Γ+Γ,
9.83 3.85 5.56 27.78 41.88 6.84 4.27
10.51 3.74 6.23 37.00 30.35 5.75 6.42
11.7841 3.7881 8.8376 49.2429 21.0434 3.6452 1.6588
Γ, Γ,+ δ δ δ+C ζ
DISCUSSION T h e X R D pattern (Figure 1 ) clearly indicates the p r e s e n c e of mainly the δ phase in the outer surface of the coating of the IF steel. T h i s is also true for the I F H S and H S Q steels. T h e cross-sectional scanning electron micrographs [Figures (2-4)] along with the corresponding X ray m a p s for Fe and Zn clearly s h o w a decrease in F e content from the substrate to the C u side. T h e gradation in Z n content is not that clear, possibly because of the overall h i g h Z n
108
Materials Processing and Texture
Effect of C o a t i n g T e x t u r e o n P o w d e r i n g Behavior of Industrially P r o d u c e d G a l v a n n e a l e d C o a t i n g
Figure 6. Φ Sections o f the O D F of the G a l v a n n e a l e d Coating o n Sample N o . i F . 2
Figure 7. Magnified V i e w of 2=0° Section of Figure 7 of S a m p l e N o . IF.
Materials Processing a n d Texture
109
Effect of C o a t i n g T e x t u r e o n P o w d e r i n g B e h a v i o r of Industrially P r o d u c e d G a l v a n n e a l e d C o a t i n g
Figure 8. Φ Sections of the O D F of the Galvannealed Coating o n S a m p l e N o . I F H S . 2
Figure 9. Magnified V i e w of Φ = 0 ° Section of Figure 8 of S a m p l e N o . I F H S . 2
110
Materials Processing and Texture
Effect o f C o a t i n g T e x t u r e o n P o w d e r i n g B e h a v i o r of Industrially P r o d u c e d G a l v a n n e a l e d C o a t i n g
F i g u r e 10. Φ 2 S e c t i o n s o f t h e O D F o f t h e G a l v a n n e a l e d C o a t i n g o n S a m p l e N o . H S Q .
F i g u r e 1 1 . M a g n i f i e d V i e w of Φ 2 = 0 ° S e c t i o n of F i g u r e 9 o f S a m p l e N o . H S Q .
Materials Processing and Texture
111
Effect of C o a t i n g T e x t u r e o n P o w d e r i n g Behavior of Industrially P r o d u c e d G a l v a n n e a l e d C o a t i n g
Table V. A v e r a g e Weigh.t L o s s (mg) during D o u b l e Olsen Test of Galvannealed Coating. IF IFHS HSQ
5.80 8.80 14.56
concentration in the coating. The percentage thicknesses of the different F e - Z n intermetallic p h a s e s present in the galvannealed coating o v e r IF, I F H S and H S Q grade steels obtained from Q D P - G D O E S profiles (Table IV) clearly indicate that the a m o u n t of δ present in the samples is m a x i m u m in the case of the IF and m i n i m u m in H S Q , the I F H S lies in between. T h e increase in the relative thicknesses of the Γ and Γι layers and the corresponding decrease in the thickness of the δ layer from IF to I F H S to H S Q can be linked to the higher a m o u n t of iron diffusion from the substrate steel t o w a r d s the coating in the a b o v e steels in t h e reverse order. This is w h a t h a s been found by the results of the wet chemical method, w h i c h m e a s u r e s the a m o u n t s of Fe present in the coating (Table III). T h e p o w d e r i n g test results corroborate the above findings pretty well. A s expected, the IF sample s h o w s the m i n i m u m weight loss, while the H S Q s h o w s the m a x i m u m weight loss (Table V ) , the value for the I F H S c o m e s in b e t w e e n . There are perceptible differences in the crystallographic textures of the coatings in the three samples. A l t h o u g h t h e major texture c o m p o n e n t s of the coatings in all the three cases are of the type {01.3} [Figures (6, 8 and 10)], a careful examination of the Φ 2 = 0 ° O D F sections [Figures ( 7 , 9 and 11)] clearly indicates that for the IF steel there exists an almost c o m p l e t e fiber of the type {01.3}; the fiber character changes into a set of discrete orientations in the samples I F H S and H S Q . B e t w e e n the latter two, the b r e a k d o w n in the continuity of this fiber is m o r e in case of H S Q than in I F H S . T h e {01.3} planes m a k e an angle - 6 0 ° with the basal planes {00.2}. Therefore, in these galvannealed samples, the basal planes are inclined at - 6 0 ° with respect to the sheet planes. observed that in electrogalvanized materials no cracking takes place w h e n the Shaffer et al . basal plane m a k e s an angle b e t w e e n 4 0 - 7 0 ° w i t h respect to t h e substrate surface. T h e presence of an almost continuous fiber of the type {01.3} as found in case of the IF sample seems to be m o s t beneficial in this respect. B r e a k d o w n of this fiber into discrete orientations appears to cause deterioration in the p o w d e r i n g resistance behavior, as exemplified by the steels I F H S and H S Q . O f course, the best p o w d e r i n g resistance observed in case of the IF sample m u s t also b e related to the least a m o u n t of Fe content in the coated layer, as c o m p a r e d to the I F H S and H S Q . 2}
CONCLUSIONS In s u m m a r y , the coating formability decreases and the p o w d e r i n g tendency increases in galvannealed coating from IF to I F H S to H S Q grade steels. It is due to the increasing degree of discontinuity of the {01.3} fiber present in the galvannealed coating textures and the increasing a m o u n t s of iron present in the coated layers from the IF to t h e I F H S to the H S Q steel.
REFERENCES ' M . M o r i s h i t a , K.K.oyama, M . M u r a s e and Y.Mori, " I m p r o v e m e n t of Corrosion Resistance of Z i n c Plated Steel by Electrodeposition of M a g n e s i u m from a M o l t e n Salt," I SI J International, 36, 6 , 7 1 4 ( 1 9 9 6 ) . C . S . L i n , M.Meshii and C.C.Cheng, "Microstructural Characterization of G a l v a n n e a l e d 2
112
Materials Processing a n d Texture
Effect of C o a t i n g T e x t u r e o n P o w d e r i n g Behavior of Industrially P r o d u c e d G a l v a n n e a l e d C o a t i n g
Coatinfgs by T E M , " ISIJ International, 3 2 , 494 (1995). S y a h b u d d i n , P . R . M u n r o e , C , S , L a k s m i a n d B . G l e e s o n , "Effect of 0.1 a n d 0.2 w t . % A l u m i n u m Addition to Z i n c o n the Inter-Diffusion b e t w e e n Z i n c and Iron at 4 0 0 ° C , " Materials Science and Engineering A , 2 5 1 , 87 (1998). C.Xhoffer, H . D i l l e n a n d B . C . D e C o o m a n , " Q u a n t i t a v e P h a s e A n a l y s i s o f G a l v a n n e a l e d Coatings b y C o u l o m e t r i c Stripping," Journal o f Electrochemistry, 2 9 , 2 0 9 (1999). ' E . T . M c D e v i t t a n d M . M e s h i i , "Electron M i c r o s c o p y Study o f the Microstructural Evolution in the Early Stages o f G a l v a n n e a l i n g " Z i n c B a s e d Steel C o a t i n g System: P r o d u c t i o n a n d Performance, P r o c e e d i n g s o f the International S y m p o s i u m , T e x a s , U S A , 127 (1998). A . B e s s e y r i a s , F.Dalard, J . J . R a m e a u and H . B a u d i n , " E l e c t r o c h e m i c a l B e h a v i o r o f Iron Z i n c Intermetallic C o m p o u n d s in a n A q u e o u s Solution Containing N a C l a n d ZnS04," Corrosion Science, 3 9 , 1 0 , 1883 (1997). A . R . M a r d e r , " T h e Metallurgy o f Z i n c Coated Steel," Progress in M a t e r i a l s Science, 4 5 , 191 (2000). V . R a n g a r a j a n , C . C . C h e n g , L . L . F r a n k s , " T e x t u r e in the ô-(FeZnio) Phase formed in Galvanneal C o a t i n g s , " Surface a n d C o a t i n g T e c h n o l o g y , 5 6 , 2 0 9 (1993). ' P . R . S e r e , J.D.Culcasi, C.I.Eisner, A.R.D.Sarli, " R e l a t i o n s h i p b e t w e e n T e x t u r e and Corrosion Resistance in Hot D i p G a l v a n i z e d Steel S h e e t s , " Surface and Coating T e c h n o l o g y , 1 2 2 , 143 (1999). H . P a r k , Z . A . S z p u n a r , " T h e Role o f T e x t u r e and M o r p h o l o g y in O p t i m i z i n g the Corrosion Resistance o f Z i n c b a s e d Electro-galvanizing C o a t i n g s , " Corrosion Science, 4 0 , 4, 525 (1998). A . W . S i l i m p e r i , D . K . M a t l o c k , J.G.Speer, "Influence o f Surface Conditioning o n the Frictional b e h a v i o r o f H o t D i p Galvanized Sheet Steel," Z i n c B a s e d Steel C o a t i n g System: Production and Performance, P r o c e e d i n g s o f the International S y m p o s i u m , T e x a s , U S A , 195 (1998). J . L . M i l l a n , "Effect o f Electrodeposited Z i n c Coating o n the Formability o f D r a w i n g Quality A l u m i n u m Killed Sheet Steel," Journal o f Materials Engineering and Performance, 1, 2 7 5 , (1992). A . C h a k r a b o r t y , R . K . R a y and D.Bhattacharjee, "Correlation b e t w e e n the Substrate Steel Texture and C o a t i n g T e x t u r e o f T w o Industrially p r o d u c e d G a l v a n i z e d a n d G a l v a n n e a l e d Steels," M a t e r i a l s Science and T e c h n o l o g y , 2 3 , 1 0 , 1 2 2 3 (2007). A . C h a k r a b o r t y and R . K . R a y , "Effect o f Substrate T e x t u r e o n the F o r m a t i o n o f ζ p h a s e in the Overlying G a l v a n n e a l e d Coatings in Industrially P r o d u c e d Interstitial Free Steels," Scripta Materialia, 5 6 , 8, 6 5 3 ( 2 0 0 7 ) . S . G a i g n a r d , C.Esling, " O p t i m i z a t i o n o f Galvanizing P r o c e s s P a r a m e t e r s to I m p r o v e the Cracking Resistance in Z i n c C o a t i n g , " I C O T O M - 1 2 , P r o c e e d i n g s o f the 1 2 International Conference o n T e x t u r e o f Materials, 2 , 9 9 7 (2001). G . V i n c e n t , C C o u n h a y e , C.Esling, " S i m u l a t i o n a n d T e m p e r Rolling of G a l v a n i z e d Steel," Solid State P h e n o m e n a , 1 0 5 , 371 (2005). A . S e m o r o z , L.Strezov, M . R a p p a z , "Orientation D o m a i n s and T e x t u r e in Hot D i p p e d G a l v a n i z e d C o a t i n g s , " Metallurgical a n d Materials Transactions A , 3 3 , 2 6 9 5 ( 2 0 0 2 ) . M . B . M o o n , J . H . S e o , Y.B.Park, " T e x t u r e and Phase Distribution in the C o a t i n g layers o f G a l v a n n e a l e d Steel S h e e t s , " Materials Science F o r u m 4 0 8 - 4 1 2 , 1209 ( 2 0 0 2 ) . " H . J . B u n g e , T e x t u r e A n a l y s i s in Materials Science, Cuvillier V e r l a g , Gottingen, 1 9 9 3 . K . P a w l i k , J.Pospiech, K . L u c k e , " T h e D e v e l o p m e n t o f a N e w Direct M e t h o d o f O D F R e p r o d u c t i o n F r o m Pole Figures a n d its Testing with the H e l p o f M o d e l Functions," P r o c e e d i n g s o f I C O T O M - 8 , Santa Fe, U S A , 105 (1987). 3
4
6
7
8
10
1 1
12
13
l4
l5
th
16
17
18
20
Materials Processing a n d Texture
113
Effect of C o a t i n g T e x t u r e o n P o w d e r i n g B e h a v i o r of Industrially P r o d u c e d G a l v a n n e a l e d C o a t i n g
'K.Pawlik, J.Pospiech, K . L u c k e , "The O D F approximation from pole figures with the aid of the A D C - m e t h o d , " P r o c e e d i n g s of I C O T O M - 9 , A v i g n o n , France, 25 (1990). 'Standard Test M e t h o d for Weight [Mass] of Coating on Iron and Steel Articles with Zinc or Zinc Alloy Coatings, A S T M Designation: A 90/A 9 0 M - 0 1 . 'S.J.Shaffer, J . W . M o r i s , H . R . W e n k , " T e x t u r e Characterization and its Applications on Z i n c Electro-galvanized Steels," Zinc Based Steel Coating S y s t e m : Metallurgy and Performance, Proceedings of the International S y m p o s i u m on Zinc B a s e d Steel Coating S y s t e m : Metallurgy and Performance, Detroit, M i c h i g a n , U S A , 129 (1990).
114
Materials Processing a n d Texture
I N C L U D I N G E L A S T I C A L A N I S O T R O P Y D U E T O T E X T U R E IN A F I N I T E E L E M E N T M O D E L FOR THE PREDICTION OF RESIDUAL STRESSES IN W E L D S K o e n D e c r o o s ' , Carsten O h m s ' , Y v a n H o u b a e r t
2
' E u r o p e a n C o m m i s s i o n , Joint Research Centre, 1755 Z G Petten, T h e Netherlands U n i v e r s i t y of Ghent, Faculty of Engineering, D e p a r t m e n t o f Metallurgy and Materials Science, 9000 Gent, B e l g i u m
2
ABSTRACT A m e t h o d is described to include welding texture in a finite element m o d e l to predict residual stresses in an austenitic weld. D u e to a preferential [100] growth in the solidification direction the weld material is transversal isotropic with the solidification direction as the fiber axis. Results s h o w a limited difference of the predicted residual stress in the w e l d w h e n considering this effect c o m p a r e d with a traditional model where elastic isotropic materials are used. 1.INTRODUCTION W e l d i n g of thick plates is an extensively used process in industry. D u r i n g a welding process the nonuniform temperature distribution and the melting/solidification processes cause thermal stress fields w h i c h lead to residual stresses. T h e k n o w l e d g e of these stresses can b e of high importance for the study of brittle fractures, stress corrosion cracking and fatigue. Finite E l e m e n t (F.E.) M o d e l i n g of t h e r m o - m e c h a n i c a l behaviour found its origin in the 1970s and is n o w a widely used tool for the determination of residual stresses caused by w e l d i n g ' . 4
1
Previous research stated that a strong texture can be formed in w e l d s , specifically strong [100] fiber textures in the solidification direction The aim of the present study is to incorporate the mechanical anisotropy caused by this texture in a finite element m o d e l and to investigate the effect o n the predicted residual stresses by c o m p a r i n g the results with an analogous m o d e l w h e r e isotropic parameters are used. 5 , 6
2.MODEL 2.1 Analysis P r o c e d u r e The physics of the t h e r m o m e c h a n i c a l welding process is described b y a w e a k l y coupled set of thermo-elastoplastic differential equations w h e n neglecting the dissipation of m e c h a n i c a l energy. The thermal p r o b l e m is solved separately from the m e c h a n i c a l one and the recorded thermal history is u s e d as load input for the mechanical p r o b l e m . T h e adding of n e w weld filler material is simulated by the " d e a t h and birth of elements t e c h n i q u e " in both the thermal and the mechanical analysis . T h e time of activation of an element is derived from its location within the w e l d in correlation with the welding speed. T h e source of heat input in the system is a variation of G o l d a k ' s ellipsoidal heat source m o d e l presented by Sabapathy . Hereby the heat source is restricted to the w e l d material and does not penetrate the base material. 15
2
16
115
Including Elastical A n i s o t r o p y D u e t o T e x t u r e in a Finite E l e m e n t M o d e l
2.2 Mechanical Finite E l e m e n t implementation The expression o f the principle of virtual w o r k is integrated over space and is, for e x a m p l e , . T h e solution within an element is written as a function of the n o d a l solutions described in b y m e a n s of shape functions. Without g o i n g into detail the solution procedure o f a load step consists o f an iterative p r o c e d u r e where for e a c h substep i a linear set of equations is solved for the incremental nodal displacement {dU,}. 1 4 1 5
[K] {du,} = {F; -r}
(i)
m
p
After each substep the total nodal displacement and the internal forces are updated. T h e iteration stops w h e n external and internal forces c o n v e r g e sufficiently. The tangent stiffness matrix [K] is thereby given by p
[K\ = p
\B' D"BdV ('
= £ elements
\B'D"BdV
( 2 )
y
p
D contains the elasto-plastic material properties, including t h e history d e p e n d i n g yield stress updating according to the material hardening law. Β is only d e p e n d i n g on the element shape functions. T h e space integral is calculated b y s u m m a t i o n of the integral over the elements w h i c h allows defining different material properties for every element without c h a n g i n g the solution p r o c e d u r e and respecting the connectivity b e t w e e n elements with different material properties by their c o m m o n n o d e s . T h e commercially available code A N S Y S has been used for these analyses facilitating the implementation of this approach with a user subroutine. 2.3 Specimen
Figure 1 : Specimen T h e s p e c i m e n used for the simulations consists of 2 plates o f 10 m m thickness w h i c h are welded together o v e r a length of 100 m m . A real s p e c i m e n is not yet fabricated at the m o m e n t . T h e w e l d is prepared in form of a V - g r o o v e with an opening angle o f 45 degrees and consists of one bead. H e n c e , t h e plates are 100 m m w i d e at the bottom and 9 3 m m at the t o p .
116
Materials Processing and Texture
I n c l u d i n g Elastical A n i s o t r o p y D u e t o T e x t u r e in a Finite E l e m e n t M o d e l
T h e m o d e l simulates a welding process with a welding speed of 15 c m / m i n and a heat input of 3.5 M J / m . Both the w e l d and the base material are from austenitic steel and their isotropic t h e r m o m e c h a n i c a l properties are taken from a study of Katsareas 1 7
3. I N T R O D U C I N G E L A S T I C A N I S O T R O P Y IN T H E M O D E L 3.1 B a c k g r o u n d 8
Because an austenitic steel single crystal is elastically anisotropic a preferential [100] orientation of the crystallites will lead to macroscopic anisotropy of a polycrystalline aggregate. For [100] fiber textures, the orientation distribution function ( O D F ) , w h i c h gives the relative v o l u m e fraction for every orientation g, is only d e p e n d i n g of the angle φ, namely the angle b e t w e e n the [100] direction and the fiber axis . 10
Austenitic crystals have cubic symmetry, t h u s 3 independent constants describe the single crystal elastic behaviour ' . T h e theory of Voigt states that with a given O D F the Stiffness tensor of a polycrystalline material can b e determined by v o l u m e averaging of t h e single crystal stiffnesses, w e i g h t e d by the O D F ' " : 7
Cl„
= \\\C {g)f{g)dg
(3)
mnpq
The theory of R e u s s states analogically that the C o m p l i a n c e tensor S can b e found by averaging the single crystals c o m p l i a n c e tensors : H i l l s h o w e d that the former theories form upper and lower bounds for the real stiffness and compliance and as a c o n s e q u e n c e the arithmetic average of b o t h forms a good estimate: 1 2
1 3
S»=i[s* (C'')-']
(4)
+
N o t e that according t o Hill the inverse of the stiffness d o e s not equal the compliance. Because of the direct relation w i t h the Y o u n g ' s m o d u l i and Poisson's ratios used as input p a r a m e t e r s in the F.E. code the second formula is used. Being transversal isotropic a material with a [100] fiber texture h a s 5 independent elastic constants. By v o l u m e averaging t h e elastic constants can be calculated as a function of the 3 independent cubic crystal constants and 2 invariants A and A ' ' . With the z-axis being t h e fiber axis t h e Voigt m o d e l i n g s c h e m e results in:
Materials Processing and Texture
·
117
Including Elastical A n i s o t r o p y D u e t o T e x t u r e in a Finite E l e m e n t M o d e l
= C ; - | ( C „ - C , ) ( l - ^ ) ( ^ - l ) J
i
2
{c ) = c ; + i ( c „ - c ) ( i - ^ x ^ - i ) 44
4
l2
2
J
;
2
A'=2
c;, =c„
-|(c„
-c,
c;
+ I(c„
-c
2
C«
=c
l2
2
-2C
4 4
)
12
-2C
4 4
)
4 4
Γ
- Γ
A = ^q *
~ (C 11 ~ C | )
7 f(2sin <0—sin 0)/()sinffc}V 2
q=
2
(6)
4
4
J
Hereby A' is only d e p e n d i n g on t h e single crystal anisotropy in a n d A only d e p e n d i n g on the distribution o f t h e fiber texture. L o o k i n g t o t h e parameter A' it c a n b e seen that with the s a m e texture t h e m a c r o s c o p i c anisotropy is higher within an austenitic steel than within a ferritic steel because of t h e greater single crystal anisotropy o f austenite . It is also t o b e noted that a perfect fiber texture, i.e. all crystals h a v i n g the [100] direction perpendicular t o the fiber axis, 4 c o r r e s p o n d s t o q = 0 and that t h e isotropic case corresponds to q = - . 8
N o t e that the R e u s s m o d e l i n g s c h e m e , n a m e l y v o l u m e averaging of t h e c o m p l i a n c e s , gives analogous results. In this study t h e arithmetic average o f t h e b o u n d s p r o v i d e d by V o i g t and R e u s s are taken. Experimental w o r k on d e t e r m i n i n g the 5 elastic p a r a m e t e r s in w e l d s h a s b e e n d o n e for e x a m p l e in ' . In practice, for a g i v e n t y p e of austenitic steel the single crystal elastic p a r a m e t e r s , unlike the isotropic bulk elastic p a r a m e t e r s , are not very well d o c u m e n t e d in a w i d e t e m p e r a t u r e r a n g e . Therefore the following a p p r o a c h is used to obtain an estimate. W e look for the single crystalline material that h a s -according t o t h e Voigt-Reuss-Hill m o d e l i n g s c h e m e - exactly t h e s a m e isotropic bulk constants Ε a n d ν a s g i v e n for a specific steel o f interest. A s a n extra condition w e i m p o s e that for every t e m p e r a t u r e the crystal anisotropy factor 5
6
9
A =2
S
n
~
S
l
2
(7)
•^44
8
for that crystal is equal to t h e o n e w e obtain for a n austenitic single crystal published in , following t h e reasoning of that austenitic monocrystal elastic parameters are relatively 4
118
Materials Processing a n d Texture
Including Elastical A n i s o t r o p y D u e t o T e x t u r e in a Finite E l e m e n t M o d e l
independent o f their material's c h e m i c a l composition " . By doing this w e obtain at every t e m p e r a t u r e Ta set o f 3 non-linear e q u a t i o n s for the 3 independent single crystal values ί , , , ί , ί , leading to a unique solution: 1 2
4 4
E(T)
2
• - - ^ =τ( 2
5
ί
7
7
>»(T)->»(T)
=
u
Γ
u
+
5ί(ί, (Γ),ί (Γ),5 (Γ))) ι
12
44
A r n
H e r e b y is, for e x a m p l e , Sl (s ,s ,s ) t
ί
'2( ..( ').%( ')' «( ))
t h e value of t h e c o m p l i a n c e tensor calculated with the
iA
m e t h o d o f V o i g t as a function o f the u n k n o w n single crystal values. A is t h e anisotropy value 8
found for a n austenitic m o n o - c r y s t a l i n . 3.2 I m p l e m e n t a t i o n in the finite e l e m e n t c o d e The solidification direction cannot b e considered a s uniform t h r o u g h o u t t h e weld. Therefore it is necessary t o calculate the solidification direction for every single e l e m e n t in the fusion zone. A s previously m e n t i o n e d t h e nature o f finite element m o d e l i n g allows every element to h a v e different material properties. T h e e l e m e n t s in t h e m o d e l are s u p p o s e d to b e small enough t o obtain a n accurate solution and b i g e n o u g h to incorporate a sufficient a m o u n t of grains in order to justify t h e averaging p r o c e d u r e s followed. D u e to insufficient k n o w l e d g e of the possible gradient o f the texture, a uniformly textured material is u s e d for every element, but t h e orientation is different according t o t h e different solidification directions. This is i m p l e m e n t e d in A N S Y S b y a priori defining t h e e l e m e n t as orthotropic. T h e solidification direction is defined as the opposite direction o f t h e heat flux o n t h e m o m e n t the a v e r a g e e l e m e n t t e m p e r a t u r e g o e s b e l o w t h e solidification t e m p e r a t u r e . It is calculated by a user subroutine after postprocessing the results o f t h e thermal run. T h e direction is given by 2 Euler angles and will determine t h e individual coordinate system o f the element in t h e m e c h a n i c a l r u n . E l e m e n t s o f the w e l d are b o r n with their orthotropic p a r a m e t e r s , elements that are m o l t e n during t h e process are initially considered a s isotropic base material and are reborn as anisotropic e l e m e n t s . 4. R E S U L T S 3 cases are studied. E a c h m o d e l h a s the exact s a m e thermal run a n d t h e s a m e mono-crystals but have a different texture a n d consequently different elastic p a r a m e t e r s . T h e cases are distinguished by the previously introduced parameter q. T h e cases q = 0 (perfect fiber texture), q=2/5 and q = 4 / 5 (isotropic case) a r e studied. A c c o r d i n g t o the m e t h o d described their elastic parameters w e r e calculated. I n F i g . 2 Y o u n g ' s m o d u l u s at r o o m t e m p e r a t u r e is plotted as a function o f the angle with the fiber axis, w h i c h is t h e horizontal axis in t h e figure. T h e solid line is t h e perfect fiber texture, t h e dashed line the case w h e r e q = 2 / 5 and t h e thick dash dotted line represents t h e isotropic case.
Materials Processing a n d Texture
119
Including Elastical A n i s o t r o p y D u e t o T e x t u r e in a Finite E l e m e n t M o d e l
Figure 2: angular variation o f Y o u n g ' s M o d u l u s with respect to the fiber axis T h e isotropic case has a direction independent m o d u l u s of 170 G P a . For the case of a perfect fiber texture a m a x i m u m m o d u l u s o f 222 G P a is found in a direction which m a k e s an angle o f 50.5° with the fiber axis. T h e minimal m o d u l u s of 86 G P a is reached a l o n g the fiber axis. In the case q=2/5 the highest value o f the m o d u l u s is 197 G P a in a direction 49.7 degrees from the fiber axis, the m i n i m u m value is, again along the fiber axis, 122 G P a . A user subroutine calculates the solidification direction for every element in the fusion z o n e . T h e results for a slice in the weld are s h o w n in Fig. 3 b e l o w from 2 different perspectives. T h e a r r o w indicates the welding direction.
Figure 3 : solidification direction for a cut in the middle of the w e l d It can be seen that the m a i n solidification direction in the w e l d tends t o align with the welding direction, except near the b a s e material and near the free surface. For all 3 cases the longitudinal stress plot for a longitudinal line at mid-thickness of the specimen is plotted in figure 4.
Figure 4 : stress in z-direction along the w e l d in the m i d d l e of the weld
120
·
Materials Processing and Texture
I n c l u d i n g Elastical A n i s o t r o p y D u e t o T e x t u r e in a Finite E l e m e n t M o d e l
The highest stresses are found in the isotropic case. T h i s is probably d u e to the fact that the zdirection is the m a i n solidification direction w h i c h the crystallographic [100]-direction, as can be seen o n t h e directional m o d u l u s plot in the figure a b o v e , tends to align with. T h e [100] direction is the direction in w h i c h an austenitic mono-crystal h a s the lowest Y o u n g ' s M o d u l u s . The stresses in the x-direction for the s a m e line in the w e l d are given in Fig. 5. N o difference b e t w e e n the m o d e l s can b e seen here.
Figure 5: stress in x-direction along the weld in the m i d d l e of the weld T h e longitudinal stresses at mid thickness on a line in the x-direction are plotted in Fig. 6:
F i g u r e 6: stress in z-direction a l o n g a line transversal to the weld at m i d thickness Again there is a small difference in the middle of the w e l d only and virtually no difference is seen in the base material. 5.CONCLUSIONS A method to incorporate the elastic anisotropy due to texture in m o d e l i n g o f residual stresses in austenitic w e l d s is presented. 3 different cases have been studied, only differing by their elastic anisotropy d u e t o different degrees of texture in the fusion z o n e . O n e of these is the 1 0 0 % isotropic case c o m m o n l y used in F E m o d e l s . Only small differences in predicted stresses are found, and only in the weld in the direction of welding. In the base material no differences are found. In that direction a difference of 35 G P a is found b e t w e e n the isotropic case and the perfect fiber texture case. This difference c a n n o t b e easily confirmed or disproved b y experimental means. T h i s study only focused on elastic anisotropy, further research to the effect o f taking into account plastic anisotropy is needed.
Materials Processing and Texture
·
121
Including Elastical A n i s o t r o p y D u e t o T e x t u r e in a Finite E l e m e n t M o d e l
6. R E F E R E N C E S 1
G. B o u c h e , J. L. B é c h a d e , M . H. M a t h o n , L. Allais, A. F . G o u r g u e s a n d L. N a z é , T e x t u r e o f w e l d e d j o i n t s o f 316L stainless steel, multiscale orientation of a weld metal deposit, Journal of Nuclear Materials, 2 7 7 , 9 1 - 9 8 ( 2 0 0 0 ) . 2
J.Goldak, A . Chakravarti, and M . Bibby, a n e w finite element m o d e l for w e l d i n g heat sources, Metall. Trans. Β., 1 5 B , 2 9 9 - 3 0 5 (1984). Η-A Kuhn, H - G Sockel, Elastic Properties o f Textured and Directionally Solidified NickelA 1 1 2 , 117-126, based Superalloys B e t w e e n 25 a n d 1200 C, Materials Science and Engineering, (1989) 3
4
W. V a n d e r m e u l e n , M . Scibetta, A . Lenaerts, J. S c h u u r m a n s , R. G é r a r d , M e a s u r e m e n t of the Y o u n g m o d u l u s anisotropy of a reactor pressure vessel cladding, Journal of Nuclear materials, 3 7 2 , 2 4 9 - 2 5 5 (2008). 5
H. M . Ledbetter, Monocrystal elastic constants in the ultrasonic study o f w e l d s , Ultrasonics, 2 3 , 9-13 (1985). B.R. D e w e y , L. Adler, R . T . K i n g , K.V. C o c k , M e a s u r e m e n t s o f Anisotropic Elastic Constants o f T y p e 3 0 8 Stainless-steel Electroslag W e l d s , Exp. Mechanics , 4 2 0 - 4 2 6 (1977). U. K o c k s , C T o m e , H . W e n k , T e x t u r e a n d anisotropy: preferred orientations in polycrystals and their affect o n materials properties. Cambridge: Cambridge University Press ( 1998). G. S y m m o n s , H. W a n g , Single Crystal Elastic Constants a n d Calculated A g g r e g a t e Properties: a H a n d b o o k , M.I.T. Pres, Boston (1971). J.F. N y e , Physical Properties o f Crystals-Their Representation by Tensors and Matrices, Clarendon, Oxford ( 1985). H.J. B u n g e , T e x t u r e Analysis in Material Science, Butterworths, London (1982). W . Voigt, L e h r b u c h der Kristallphysik, Teubner, Leibzig (1928). A. Reuss, B e r e c h n u n g d e r fliessgrenze von mischkristallen auf grund der plastizitötsbedingung fur einkristalle, Z. Angew. Math. Mech., 9 , 4 9 - 5 8 (1929). R. Hill, T h e elastic behaviour o f a crystalline aggregate, Proc. Phys. Soc. London, A 6 5 , 349 (1952). L E Lindgren, numerical m o d e l l i n g o f w e l d i n g , Comput. Methods Appl. Mech. Engrg, 1 9 5 , 6 7 1 0 - 6 7 3 6 (2006). A N S Y S , A N S Y S 11.0 C o m p l e t e user's M a n u a l set P.N. Sabapathy, M . A . W a h a b , M.J. Painter, N u m e r i c a l m o d e l s o f in-service welding o f gas pipelines, International conference on Advances in Materials and Processing technologies, (1999). 6
7
8
9
1 0
1 1
1 2
1 3
1 4
1 5
1 6
1 7
D.E. Katsareas, A . G . Y o u t s o s , Residual Stress Prediction in Dissimilar Metal W e l d Pipe Joints using the Finite E l e m e n t M e t h o d , Materials Science Forum, 4 9 0 - 4 9 1 , 53-61 (2005). R.A. A d a m e s k u , E.A. Mitjushov, L.L. Mitjushova and K . S . J u s h k o v , Invarianten der Anisotropie elastischer Eigenschaften v o n texturierten kubischen Metallen, Z. Metallkd., 16, 747 (1985). 1 8
1 9
Metals H a n d b o o k , Properties and Selection: Irons, steels and H i g h Performance alloys, tenth ed., vol 1, ASM International (1990).
122
Materials Processing a n d Texture
T E X T U R E A N D M I C R O S T R U C T U R E E V O L U T I O N IN A Fe-Si C G O S H E E T D U R I N G T H E PROCESSING ROUTE BEFORE SECONDARY RECRYSTALLIZATION. 1
2
3
2
3
1
Francisco C r u z - G a n d a r i l l a , Richard P e n e l l e ' , Thierry B a u d i n ' , Hector M e n d o z a L e o n and J. G. Cabafias-Moreno Instituto Politécnico N a c i o n a l , E S F M , D e p t o . de Ciencia d e Materiales, U n i d a d Prof. A L M , Edif. 9, Z a c a t e n c o , M e x i c o , D.F., 0 7 3 3 8 , M e x i c o . C N R S , U M R 8 1 8 2 , I C M M O , Laboratoire d e Physico-Chimie de l'Etat Solide, O r s a y , F - 9 1 4 0 5 ; University o f Paris-Sud, Orsay, F - 9 1 4 0 5 . 1
1
2
3
ABSTRACT. T h e present w o r k deals with the origin of the G o s s grains present in the primary recrystallized structure. Consequently, the generation of the m a i n texture c o m p o n e n t s and their associated microstructures w a s followed from the hot rolled to the primary recrystallized state at each step of the processing route. T h e texture w a s quantitatively characterized by X-ray diffraction and E B S D . In the hot rolled state, the microstructure is m a d e up o f b a n d s of subgrains or recrystallized grains. N e a r the surface and a quarter o f the sheet thickness ( Q S T ) the texture possesses a strong G o s s c o m p o n e n t , its strength decreasing w i t h increasing depth d o w n t o the center. Conversely, the α fiber texture c o m p o n e n t s increase in intensity t o w a r d s the center of the sheet. After the first stage of cold rolling, the strength of the Goss orientation drops d o w n to the benefit of the γ texture c o m p o n e n t s , (111)[1-21] and ( 1 1 1 ) [ - 1 - 1 2 ] . F r o m E B S D determinations, it is found that the G o s s grains rotate mainly towards these orientations, the volume fractions of which b e c o m e approximately equal to 15 % each. T h e α texture c o m p o n e n t s are the major components at the center of the sample and do not significantly change until the last annealing treatment, w h e n their v o l u m e fractions decrease noticeably. T h e s e c o n d cold rolling step results in the stabilization of the (111)[1-21] and ( 1 1 1 ) [ - 1 12] orientations formed close to Q S T . Grains having these γ texture orientations also s h o w shear bands inclined 35-40° relative to the rolling plane. A n interesting feature observed in the shear bands located in grains w i t h t h e (11 l ) [ l - 2 1 ] orientation is the formation of grains w i t h either a (111)[1-10] orientation or the G o s s orientation. Conversely, grains with a (111)[0-11] orientation appear only o n shear bands located within grains with a (111)[1-21] orientation. It w a s also noticed that s o m e G o s s grains w e r e p r e s e n t b e t w e e n lamellar grains too. W i t h the last annealing, the strength of the α fiber decreases to the benefit of the (111 ) [1 21] and (111) [1-12] orientations and the r a n d o m c o m p o n e n t , while the v o l u m e fraction of G o s s grains b e c o m e s negligible. INTRODUCTION A l t h o u g h industrial steelmakers h a v e the k n o w h o w to p r o d u c e the G o s s texture in Fe3%Si s t e e l s " , the origin and d e v e l o p m e n t o f the G o s s nuclei throughout the processing route are still not well understood. In previous w o r k s , w e have studied the texture and microstructure evolution in a C G O F e - 3 % S i hot rolled sheet at different stages o f a processing r o u t e similar to the one applied in industry ' ' . In the hot rolled state, the microstructure is heterogeneous and consists of small and large grains, the latter forming so-called Morphological B a n d s ( M B s ) . These b a n d s are parallel to the rolling direction and their m a i n orientations h a v e b e e n determined as (001)[110] and (112)[1-10] ( a fiber), { l l l ) < l l - 2 > , {111}<0-11> (γ fiber), and (110)[001] 1
2
3
4
5
123
T e x t u r e a n d M i c r o s c t r u c t u r e Evolution in a F e - S i C G O S h e e t D u r i n g P r o c e s s i n g
(Goss). H o w e v e r , the v o l u m e fractions of these c o m p o n e n t s vary w i t h the specific treatment and across the sheet thickness, as determined from the reconstructed E B S D images and D R X pole figures. However, the experimental pole figures obtained by X R D s h o w important texture differences depending o n the actual location in the samples, that is, from the surface, Q S T or center. Moreover, the X R D data d o not exhibit the m m m symmetry generally observed by neutron diffraction after hot rolling. This means that the texture c o m p o n e n t s are not randomly distributed in t h e v o l u m e of the sheet samples, although there exists some s y m m e t r y relative to a plane defined by the normal and rolling directions. The present w o r k is focused o n the evolution of the v o l u m e fractions o f the different texture c o m p o n e n t s at different stages of the processing route, p a y i n g special attention to the formation of the t w o equivalent orientations (111)[1-21] and (111)[-1-12]. MATERIALS AND EXPERIMENTAL PROCEDURE The hot rolled sheet used for this study w a s provided by Usinor. Its chemical composition is given in Table 1. C
Mn
S
Si
Cu
0.04
0.06
0.02
3.2
0.16
S a m p l e s parallel t o the rolling plane were prepared from locations at the surface, at a quarter o f the sheet thickness ( Q S T ) and at half the sheet thickness (center) for p o l e figure determination by X-ray diffraction ( X R D ) . Samples perpendicular to the rolling ( R D ) and transverse ( T D ) directions w e r e prepared for microstructure characterization by E B S D . Five different thermo-mechanical treatments w e r e applied; they are referred as A (hot-rolled), Β ( A + 7 0 % cold-rolling), C (B + 4 0 s anneal at 8 5 0 ° C under H atmosphere), D (C + 6 0 % cold-rolling) and Ε ( D + 4 m i n anneal at 830°C under H atmosphere). Global textures w e r e determined by X ray diffraction ( C o - Κα radiation) using a L ü c k e texture goniometer equipped with an 1NEL position sensitive detector ( C P S 120). T h e Orientation Distribution Functions ( O D F ) w e r e calculated by the A D C m e t h o d ( L A B O T E X 3.0 software) from the measured p o l e figures. T h e v o l u m e fractions of the m a i n texture c o m p o n e n t s w e r e calculated using the M o d e l Functions Method. Local texture w a s characterized by E B S D ( T S L - E D A X system installed on a X L 3 0 S F E G / S I R I O N S E M from F E I C o m p a n y ) . The orientation m a p s and volume fractions of t h e texture c o m p o n e n t s w e r e determined with the orientation i m a g i n g microscopy ( O I M * ) software. 2
2
RESULTS AND DISCUSION Fig. 1 s h o w s the O I M * m a p s of the w h o l e longitudinal plane defined by the R D and the normal direction (ND) after the different treatments. The most relevant feature for all conditions is the presence of Morphological B a n d s ( M B s ) of about 10 μηι thickness along the N D direction. A s a consequence, there are abrupt textures changes w h e n crossing through the M B s along the N D direction. M o r e specifically, •
124
In the hot rolled state (treatment A ) , the orientation of the M B s varies t h r o u g h the sheet thickness. Other important features at this stage are the presence of a strong Goss
Materials Processing a n d Texture
T e x t u r e a n d M i c r o s c t r u c t u r e Evolution in a F e - S i C G O S h e e t D u r i n g P r o c e s s i n g
c o m p o n e n t at Q S T locations and a strong α fiber at the center; there are also a few Goss grains at the center. • After t h e first rolling step (treatment B), the G o s s c o m p o n e n t rotates t o w a r d s γ fiber orientations, but the b a n d structure is retained. • After t h e first annealing step (treatment C), the M B structure is still present, but s o m e orientations change; h o w e v e r , the α fiber ( w h i c h is in principle a stable orientation under rolling deformation) r e m a i n s at the center position. • After the s e c o n d cold rolling step (treatment D ) , as for treatment B , s o m e orientations have u n d e r g o n e s o m e c h a n g e s . • After the last annealing (treatment E), the recrystallized structure mostly consists of grains of r a n d o m orientation and s o m e G o s s grains at Q S T locations. At center locations, the grains are visibly elongated. Figures 2 and 3 s h o w s t h e v o l u m e fractions of the main texture c o m p o n e n t s at each step of t h e processing route, as determined by X R D for center and surface locations of the samples. T h e results s h o w that the v o l u m e fractions of the (111)[1-21] and (111)[-1-12] c o m p o n e n t s at the center increase to about 1 5 % in going from the hot rolling state to the last
Materials Processing and Texture
·
125
T e x t u r e a n d M i c r o s c t r u c t u r e Evolution in a F e - S i C G O S h e e t During P r o c e s s i n g
annealing stage, although a temporary d r o p is observed after the second cold rolling step. T h e corresponding v o l u m e fractions of the (111)[011] and (111)[1-10] c o m p o n e n t s are negligible. T h e a m o u n t of e a c h of the α c o m p o n e n t s (112)[1-10] and (113)[1-10] r e m a i n s constant, w h e r e a s the v o l u m e fraction of the G o s s c o m p o n e n t (001)[1-10] gradually decreases u p to the last annealing. T h e (110)[001 ] G o s s c o m p o n e n t is practically nonexistent.
Fig. 2: Evolution of the v o l u m e fraction of the main texture c o m p o n e n t s at the center of the s a m p l e s throughout the processing route.
Fig. 3 : Evolution of the v o l u m e fractions of t h e m a i n texture c o m p o n e n t s at the surface of the samples throughout the processing route. A s determined by X R D at Q S T locations (not s h o w n ) , the v o l u m e fractions o f the pairs of equivalent orientations (111)[1-21] + (111)[-1-12], and (111)[011] + ( 1 1 1 ) [ 1 - 1 0 ] , increase from the hot rolling state to the last annealing stage, reaching values of about 1 0 % and 5 % , respectively. F o r the α c o m p o n e n t s (112)[1-10] and ( 1 1 3 ) [ 1 - 1 0 ] , the v o l u m e fractions remain l o w and stable (about 5%) until the last annealing step, w h e r e a s that of the (001)[1-10] orientation continuously increases u p to 3 0 % until the step C is reached, subsequently falling to a v a l u e of about 2 0 % after step E. The G o s s v o l u m e fraction ( 3 5 % ) is very h i g h at a Q T S , but it drops d o w n during step Β and r e m a i n s very l o w till the end of the processing route.
126
Materials Processing and Texture
T e x t u r e a n d M i c r o s c t r u c t u r e Evolution in a F e - S i C G O S h e e t D u r i n g P r o c e s s i n g
T h e v o l u m e fractions calculated from E B S D data confirm those obtained b y X R D ; however, erratic variations are observed because of the microstructural heterogeneity related to the presence of t h e M B s . After the first cold rolling step, the equivalent orientations (111)[1-21] and (111)[-1-12] are present through the sheet thickness, as indicated in Fig. 4. This figure also s h o w s that they are not randomly distributed. A s a matter of fact, o n the right side of Fig. 4 the v o l u m e fraction of the (111 ) [-1 -12] orientation is larger than that of its equivalent orientation, (111 ) [ 1 -21 ] , w h e r e a s the contrary is true on the left side of the figure (the color code is the s a m e as in Fig. 1 ).
RD
ND Fig 4: E B S D m a p s after the first cold rolling step (stage B), plane N D - R D .
T h e same kind o f a s y m m e t r y is o b s e r v e d in the pole figures d e t e r m i n e d by X D R from center locations of different hot rolled samples, Fig. 5. T h e p o l e figures in Figs 5a and 5b correspond to the case in w h i c h the v o l u m e fractions of the two symmetrical orientations, (111) [1-12] and ( 1 1 1 ) [ 1 - 2 1 ] , are different. O n the contrary, the orthotropic s y m m e t r y observed in the pole figure o f Fig. 5c, w h i c h w a s d e t e r m i n e d by n e u t r o n diffraction, is a c o n s e q u e n c e of the averaging of m a n y M B s having either o n e of the orientations (111) [-1-12] or (111)[1-21] and others.
a)
b)
c)
Fig 5: {110} pole figures. T h e pole figures of the hot rolling state (a and b ) w e r e determined by X R D from the center of the sample (the triclinic symmetry of the texture is observed). The pole figure in (c) w a s d e t e r m i n e d b y n e u t r o n diffraction for the s a m e hot rolled sample. In this case, the orthorhombic s y m m e t r y of the texture is statistically preserved. After the first cold rolling step, the reconstructed microstructures s h o w that the G o s s grains w h i c h are e m b e d d e d in the polycrystalline matrix rotate mainly t o w a r d s o n e of the t w o for equivalent { 1 1 1 } < 1 1 2 > orientations, unlike the case o f a cold rolled, G o s s single c r y s t a l which the t w o symmetrical c o m p o n e n t s (111)[1-21] and (111)[-1-12] w e r e simultaneously 6-7
Materials Processing a n d Texture
·
127
T e x t u r e a n d M i c r o s c t r u c t u r e E v o l u t i o n in a F e - S i C G O S h e e t D u r i n g P r o c e s s i n g
observed. Furthermore, s o m e remaining G o s s grains are o b s e r v e d b e t w e e n t h e lamellar grains in the M B s . A subsequent annealing (step C ) d o e s not fundamentally modify t h e relative strength of the different texture c o m p o n e n t s . H o w e v e r , after the s e c o n d cold rolling step, the fine structure in the regions with {111J < 1 1 2 > orientations s h o w narrow b a n d s c o r r e s p o n d i n g to t h e t w o equivalent orientations (111)[-1-12] and (111)[1-21], see Fig. 6 (the color c o d e is the s a m e as in Fig. 1). F i g u r e 6a displays subgrains with a (111)[1-10] orientation and shears b a n d s ' e m b e d d e d in a (111)[-1-12] M B , w h e r e a s Fig. 6 b s h o w s subgrains with (111)[0-11] orientation and shear b a n d s e m b e d d e d in a (111)[1-21] M B together with (111)[1-10] grains elongated along the rolling plane. T h e (111)[1-10] orientation is related to (111)[-1-12] by a rotation of 30° about the [111]; similarly, the (111)[0-11] orientation is related to (111)[1-21] b y a30° rotation around [111]. T h e G o s s orientation is also found within the shear b a n d s in the M B ' s with (111)[1-21] o r i e n t a t i o n , Fig. 6b.
a) b) F i g 6: E B S D m a p s of t w o Morphological B a n d s , (a) (111)[-1-12] and (111)[1-10] orientations after treatment B . (b) (111)[1-21] and (111)[0-11] orientations after treatment D .
After the last annealing treatment, primary recrystallization is almost c o m p l e t e . In order t o determine t h e orientation of the first n e w grains during the early stages of this processing route, interrupted annealings w e r e performed after step D at 7 0 0 ° C during 1 and 3 minutes. After annealing for 1 min at 7 0 0 ° C , t h e microstructure is mainly recovered, with only a few n e w recrystallized grains visible under the S E M . Figure 7 clearly exhibits the existence of t w o n e w G o s s grains within a prior M B . T h e s e n e w G o s s grains w e r e probably present as subgrains in the shear b a n d s observed in a (111)[-1-12] oriented cold rolled grain. O n the other h a n d , F i g s . 8b and 8d s h o w the p r e s e n c e of n e w grains with ( 1 1 1 ) [ 1 - 2 1 ] , and (111)[-1-12] orientations inside prior {111} < 1 1 2 - 1 1 0 > M B ' s after 3 min annealing. Therefore, the retention of the {111}<112> deformation texture can be explained by either in-situ recrystallization or strain induced boundary migration (S1BM), since b o t h n e w orientations are already present in the deformed state within the M B s . T h e n e w G o s s grains are obviously present in the regions with γ texture.
128
Materials Processing and Texture
T e x t u r e a n d M i c r o s c t r u c t u r e Evolution in a F e - S i C G O S h e e t D u r i n g P r o c e s s i n g
T o s u m m a r i z e , texture and microstructure in the M B s are h e t e r o g e n e o u s across the sheet thickness; they also vary throughout the processing route. T h i s feature is particularly o b v i o u s inside the (111)[1-21] and (111)[-1-12] M B ' s .
Fig 7: G o s s grains after 1 m i n annealing at 7 0 0 ° C . <111)[1-21]
(111)[0-11]
(111)[-1-12] b) Fig 8: E B S D m a p s (a) after 1 and (b) 3 m i n at 700°C. E B S D m a p s reconstructed from c o m p o n e n t contrast after (c) 1 m i n and (d) 3 m i n at 7 0 0 ° C (the color code is t h e s a m e as figure 1). a)
CONCLUSIONS T h e m a i n results of the present w o r k are: •
T h e hot rolling texture is characterized by a strong G o s s c o m p o n e n t at Q S T locations and an α texture c o m p o n e n t at the center of the sheet.
•
D u r i n g the first cold rolling step, the Goss texture transforms into the γ fiber, but contrary to t h e studies on single crystals, the G o s s grains e m b e d d e d in a polycrystalline material rotate t o w a r d s one of the t w o equivalent orientations (111)[1-21] and (111)[-1-12]. D u r i n g the second cold rolling step the deformation of the {111 }<112> grains occurs by shear b a n d i n g . T h e orientation of t h e subgrains in t h e s e M B ' s is either (111)[-1-12] and (111)[1-10] or (111)[1-21] and (111)'[0-11]. Consequently, the (111)[1-10] and (111)[0-
Materials Processing a n d Texture
129
T e x t u r e a n d M i c r o s c t r u c t u r e Evolution in a F e - S i C G O S h e e t D u r i n g P r o c e s s i n g
11] orientations are related t o (111)[-1-12] or (111)[1-21] b y a 3 0 ° rotation about t h e [ 1 1 1 ] direction. •
In the cold rolled state, the G o s s subgrains located within s h e a r bands of the (111)[1-21] and (111)[-1-12] M B ' s g r o w during t h e last annealing t o give rise t o t h e n e w G o s s grains.
•
Simultaneously, n e w (111)[1-21] a n d (111)[-1-12] grains within prior {111} < 1 1 2 - 1 1 0 > M B ' s g r o w during subsequent annealing at 7 0 0 °C. F r o m this, t h e retention of the { 1 1 1 } < 1 1 2 > deformation c o m p o n e n t c a n b e explained either b y in-situ recrystallization or S I B M , since these orientations are clearly present in t h e M B ' s in t h e deformed state.
ACKNOWLEDGEMENTS F C G a n d G C M are recipients o f C O F A A - I P N fellowships. REFERENCES ' N . R o u a g . Influence d e la texture cristallographique et d e la spécialité d e s j o i n t s d e grains sur l'anisotropie d e migration d e s j o i n t s entourant u n grain d'orientation {110} < 0 0 1 >, a u cours des premiers stades d e la recristallisation secondaire, dans d e s tôles d e F e - 3 % s i en p r é s e n c e d e précipités d e A I N et M n S . Ph. D . thesis (Université d e Paris S u d , 1988). M . M a t s u o , ISIJ International, 2 9 , 809, (1989). F. Cruz-Gandarilla, T. B a u d i n , R. Penelle a n d H. M e n d o z a - L e o n . P r o c e e d i n g s o f t h e S e c o n d Joint International Conferences o n Recrystallization a n d Grain G r o w t h ( R e X & G G 2 , S F 2 M ) . Materials Science F o r u m 4 6 7 - 4 7 0 , 123-128 (2004). 2
3
4
F . Cruz-Gandarilla, R. Penelle, H - M . L e o n , T. Baudin a n d J. G. C a b a n a s - M o r e n o . A study o f
local microstructure a n d texture heterogeneities in a C G O F e 3 % S i alloy from h o t rolling t o primary recrystallization. Materials Science F o r u m 4 9 5 - 4 9 7 , 4 8 3 - 4 8 8 , (2005). F. Cruz-Gandarilla, T. B a u d i n , M . H . M a t h o n , R. Penelle, H . M e n d o z a - L e o n a n d J . G . C a b a n a s M o r e n o . " C h a r a c t e r i z a t i o n o f global and local textures in h o t rolled C G O F e 3 % S i " . Materials Science F o r u m 5 0 9 , 25-30 (2006).. 6
D . Dorner, L . L a h n a n d S. Zaefferer. Proceedings o f the S e c o n d Joint International Conferences o n Recrystallization a n d Grain G r o w t h ( R e X & G G 2 , S F 2 M ) . Materials Science F o r u m V o l . 4 6 7 470, 129(2004). D . Dorner, L . L a h n a n d S. Zaefferer. Survival of G o s s grains d u r i n g cold rolling of a silicon steel single crystal. Materials Science F o r u m 4 9 5 - 4 9 7 , 1061-1066, (2005). B . D . Cullity E l e m e n t s o f X - R a y Diffraction. Edit. A d d i s o n - W e s l e y (1959) K. Ushioda, W . Hutchinson, ISIJ International, Vol. 2 9 , 8 6 2 (1989). 7
8
9
130
Materials Processing a n d Texture
TEM STUDY OF THE INHIBITION LAYER OF COMMERCIAL HOT-DIP GALVANIZED STEELS (2)
( 3 )
Jack E L I A E . P E T I T J - S L E C O M T E B. G A Y , V. P I T C H O N L E T A M , C N R S - U M R 7078, University of Metz, Ile du Saulcy. 57045 Metz, France U M I C O R E Research, 2250 O L E N , Belgium L M S P C , C N R S - U M R 7 5 1 5 , Université Louis Pasteur - Strasbourg, France ( 1 )
( 2 )
( 3 )
ABSTRACT: Hot-dip galvanized steels are largely used in automotive industry. The control of microstructure of the coating is mandatory to improve the mechanical performance to enhance formability, adhesion and d a m a g e resistance. In continuous galvanization process, small amounts of a l u m i n u m are added to galvanization alloys, the fast reactivity of aluminum with steel produces a thin Fe-Al layer w h i c h inhibits the formation of the brittle Fe-Zn intermetallics. Extensive research has been m a d e to understand the formation process of the inhibition layer that governs the nucleation and the growth of the coating. T h e extreme thinness of this layer, around 50 to 150 nm m a k e s the buried layer characterization delicate on plane view because the samples are difficult to prepare for T E M studies. In this paper, the inhibition layer w a s analyzed after the Zinc coating and the steel substrate have been stripped away by selective electrochemical dissolution. T E M imaging, selected area electron diffraction and EDX analysis have been made on the samples of GI galvanized steel at 0.35%A1 in the bath. The double layer model have been confirmed with big FeiAls grains spread on a continuous layer of nanograins at Fe interface, and a Zn-alloy layer surrounding the Fe2Al5 grains on the Zn interface has also been identified.
INTRODUCTION: Hot-dip Galvanizing is efficient to protect steel against corrosion. The coating is applied by immersing the steel products in a bath of molten zinc. In a pure Zinc bath, the process produces an overlay (Zinc coating), and an interfacial layer containing a series of Fe-Zn intermetallic c o m p o u n d s [1]. The reaction is affected by the bath time, temperature, composition and the steel grade. Fe-Zn c o m p o u n d s are usually brittle and harder than steel so their formation must be prevented on steel sheet to avoid the cracking and galling of the coating during the forming and stamping process. For this purpose, small amounts of A l u m i n u m are added to the Zinc bath (0.15- 0.35 wt % ) in order to inhibit the Fe-Zn reactions. A s a matter of fact, a thin Fe-Al layer from 50 to 150 n m grows and passivates the surface of steel. This layer governs the adhesion and prevents fracture between the steel substrate and the Zinc coating. Many studies have been m a d e to understand the behavior and the influence of the inhibition layer on Hot-dip Galvanized steels: Mechanical studies of the coating [2, 3, 4, 5, 6] demonstrate that the ductility increase by refining the grain size of the inhibition layer and controlling their orientation. After having been considered as a single layer in previous studies, recent studies about the morphology of the inhibition layer have reported a double-layer model for Al concentrations above 0 . 1 5 % and a single layer below 0.15%A1 [7], The double layer
131
T E M S t u d y of t h e Inhibition L a y e r of C o m m e r c i a l H o t - D i p G a l v a n i z e d S t e e l s
model describes that the inhibition layer consist of a first sub-layer with very small crystals and a second sublayer with big crystals. Other studies were devoted to the formation kinetics of inhibition layer: T h e r m o d y n a m i c studies on the ternary phase diagram Fe-Al-Zn at 450°c performed respectively by Perrot (1992), Yamaguchi (1995) and T a n g (1997) have identified the ( F e A l Z n ) as an equilibrium p h a s e with the liquid Zn [8,9,10]. The Zinc solubility in F e 2 A l w a s performed in order to understand the formation of F e 2 A l ; Z n [11]. T a n g studied the Nucleation and growth and proposed a model in which Al uptake at the interface is s h o w n to be a t w o stage process. He calculated the critical nucleus sizethe energy barrier for the nucleation and the rate of Al uptake [12]. Toussaint have proposed a m o r e accurate kinetic law based on T a n g model and experimental data [13], and recently, Giorgi have developed a n e w theorical model that consider the inhibition layer formed from Fe2Al5Zn according to the ternary diagrams to describe the kinetics of galvanizing reactions that occur at the steel/liquid zinc interface [14, 15, 16]. The model calculates the concentration profiles for iron and a l u m i n u m in the boundary diffusion layer, the nucleation rate, the average critical nucleus radius and the time necessary to form a continuous layer. 2
5
x
5
x
x
Despite this background, detailed description of topological and crystallographic structure of the inhibition layer is still lacking, because the layer is very thin and hard to isolate for T E M studies. In this paper, w e report a T E M analysis of the inhibition layer produced on industrial GI sheets with 0.35 A l - w t % in the Z n bath. The layer w a s m a d e accessible by stripping the Zn overlay and the steel substrate using selective electrochemical dissolution. The samples were analyzed by T E M imaging, diffraction and E D X . EXPERIMENT: The galvanized samples were produced under typical industrial conditions at the Arcelor Eurogal continuous galvanizing line. The used substrate w a s an Ultra-Low Carbon ( U L C ) steel gauge of 6 0 0 μ π ι , galvanized in a Zinc bath containing 0 . 3 5 % A l u m i n u m at 4 6 0 ° C bath temperature with final coating thickness around 1 Ομιη. The selective electrochemical dissolution has been m a d e from the steel face. Samples were cut into pieces making slightly more than 1 c m then mechanically ground to achieve around 70 μιη of thickness. The used Electrolyte w a s 3.2M NaCl and 0.34M Z n S Û 4 dissolved in distilled water. During the dissolution, the p H o f the Electrolyte decrease slightly, so 0.1 M H2SO4 have been used to maintain the p H around 4.5 by adding some droplets in the solution when the p H decrease. In the selective electrochemical dissolution experiments, the sample is deposited on the electrode, so w e can distinguish between the layers by m e a s u r i n g the electrode electrochemical potential. The applied current density w a s around 15 m A / c m that can achieve a dissolution velocity of 5.53nm/s for steel surface, and 7.14 nm/s for the Zinc surface. 2
2
RESULTS: After Electrochemical Dissolution, most of the sample is formed from heterogeneous surfaces, including the inhibition layer and s o m e parts o f coating or substrate. The inhibition layer can be observed directly at the borders of the holes. In this paper, four observations of the surface have been analyzed:
132
Materials Processing a n d Texture
T E M S t u d y of t h e Inhibition L a y e r of C o m m e r c i a l H o t - D i p G a l v a n i z e d S t e e l s
Global analysis: In Figure 1. Global analysis has been made o v e r the heterogeneous surface. E D X analysis shows Al rich signal. Diffraction rings over this surface are "spotty" and the indexed rings matches with F e A l (Table I). 2
3
Figure 1. Global analysis of the sample, w e can see: (a) the heterogeneous surface, (b) the diffraction rings including F e A l spots, and (c) the E D X analysis. 2
5
Double Layer: The next pictures show that the inhibition layer is constituted of t w o sublayers, the first is a thin layer with small grain size (Figure 4) where the nucleation of the F e A l j and F e A ^ grains begins, and the second sublayer contains F e A l and F e A l grains, surrounded by a Zn-alloy layer (Figure 3). Figure 2 shows the global morphology of the double layer, with the Fe2Al 2
2
5
3
5
Materials Processing a n d Texture
133
T E M S t u d y of t h e Inhibition L a y e r of C o m m e r c i a l H o t - D i p G a l v a n i z e d S t e e l s
grains (dark points). Diffraction rings s h o w that most of the spots are located on the 2 2 1 . 311 and 130 rings of F e A I (Table I). 2
;
Figure 2. T E M picture of the D o u b l e layer (a), and its diffraction ring with marks on F e A l spots (b), and Global E D X analysis (c) 2
134
Materials Processing a n d Texture
5
T E M S t u d y of t h e Inhibition L a y e r of C o m m e r c i a l H o t - D i p G a l v a n i z e d S t e e l s
Table I. Indexed patterns for Fig 1 -b and Fig 2-b, Ε = 200 K e V , λ = 0.00251 n m , C = 593.97, d = C/D (interatomic distance) and D = Diameter in pixels D (Pixels)
d-hkl(Â)-Fig.l
d-hkl (Â) - Fig.2
hkl
49.19
-
4.9172
110
57.94
3,835 3,2
3,835 -
200
69.44 69.6
3,1927
3,1927
111
87.29
2,5454
2,5454
021
Fe Al 2
5
020
93.59
2,3742
2,3742
310
104.77
2,1208
2,1208
221
105.81
2,1 2,0669
002
107.5
2,1 2,0669
108.11
2,0553
2,0553
130
115.06
1,9311
1,9311
112
115.88
1,9175 1,8461
1,9175 -
400
120.36 120.64
1,8419
-
202
126.56
1,7557
-
022
135.65
1,638
-
330
138.88
1,6
-
040
139.19
1,5964
-
222
145.08
1,5316
-
421 331
311
131
145.61
1,526
-
148.96
1,4917
1,4917
510
150.48
1,4766
1,4766
240
151.27
1,4689
1,4689
132
156.92
1,416
-
402
159.51
1,393
-
241
Zinc-alloy layer: The Zinc-alloy layer (Fe-Zn) surrounding the Fe2Al grains and the Zinc spangle traces can be easily identified in Figure 3 , this layer has an intense Zn-Κα signal. Its chemical nature has not been identified exactly yet. It is believed that this layer can be easily dissolved when the sample is attacked by acid, which might explain why it has not been reported in any previous paper. 5
Materials Processing and Texture
135
T E M S t u d y of t h e Inhibition L a y e r of C o m m e r c i a l H o t - D i p G a l v a n i z e d S t e e l s
Figure 3: Zinc-allo> layer, we can see the zinc-coating traces (a), and the F e i A l , grains (b) and E D X analysis from the center (c) and close to the border (d).
136
·
Materials Processing and Texture
T E M S t u d y of t h e Inhibition L a y e r of C o m m e r c i a l H o t - D i p G a l v a n i z e d S t e e l s
First sublayer: Figure 4 shows the first sublayer, with small nanograins of 4nm, which can explain the diffraction rings broadening (Figure 6). Diffraction rings and E D X analysis show that we have a possible mixture of FeAl, Fe Al5, and FeAlj. These phases are thermodynamically possible according to the ternary phase Diagram Fe-Al-Zn Γ81 2
Figure 4: Global Image (a) and grain picture (b) of the first sublayer E D X analysis at different points of the first sublayer (Figure 5) shows an Al/Fe atomic ratio from 0.94 to 1.7, so the global layer composition can be written in the following form: F e A l = a F e A l J + ß F e A l + yFeAl With a + β + γ =1 Nickel signals are c o m i n g from microscope grid, and are considered as "noise signals", their intensity increases if the analyzed zone is close to the grid or if the layer is very thin so the sample will not generate enough x-rays photons to neglect the noise signals. a
b
2
3
Figure 5: E D X analysis c o m i n g from two different zones of the first sublayer, Al/Fe ratio varies from 0.94 to 1.7
Materials Processing and Texture
137
T E M S t u d y of t h e Inhibition L a y e r of C o m m e r c i a l H o t - D i p G a l v a n i z e d S t e e l s
Diffracted rings from the first sublayer have been compared to the three possible diffraction data of F e A l , F e A l and FeAl (Figure 7). A S T M p o w d e r diffraction of F e A l s data can not cover all the diffraction rings that we got (limited to the dashed line), so Diffraction ring positions and intensity for Fe2Al5 have been calculated using the classical formulas given by Williams & Carter [18]: D (Pixels) = C/dhki (where C is a constant depending o n wavelength and C a m e r a length with f, the atomic scattering factors for electrons from Edington tables 2
5
3
2
Figure 6: Diffraction rings of the fist sublayer (on the Top) c o m p a r e d with FeAl, Fe2ÄU and FeAlj diffraction rings on the right side. The dashed line in the F e A U profile is the experimental limit of Fe Al5 rings according to A S T M tables. 2
2
138
Materials Processing a n d Texture
T E M S t u d y of t h e Inhibition L a y e r of C o m m e r c i a l H o t - D i p G a l v a n i z e d S t e e l s
Table II. Quantitative analysis of Diffraction rings spectra c o m i n g from the first sublayer R- Exp
51-65
Fe Al 2
FeAl
3
d(khl)
R(pix)
d(khl)
R(pix)
d(khl)
52
2.124
221
51
2.16
115
38
2.895
100
53
2.11
2
53
2.1
620
54
2.049
110
54
2.067
311, 130
53
2.08
205
57
1.938
112
54
2.06
602
54
2.05
25
54
2.04
423
55
2.02
40
56
1.99
530 66
1.673
111
77
1.448
200
86
1.295
210
1.183
211
1.489, 1.473
510, 132
75
1.49
605
78
1.417
402
75
1.48
60-8
80
1.396
241
77
1.45
245
77
1.44
642
89
1.25
408
75 74-82
FeAl
5
R(pix)
87
1.277, 1.276
89
1.245
530
91
1.22, 1.217
223, 512
94
92
1.209
313
108
1.024
220
100
1.106
351
115
0.965
300
121
0.916
310
133
0.837
222
042, 600
102
1.092
602
104
1.072, 1.069
532, 060
105
1.055
4
106
1.045
711
109
1.019
243
90-115
H i g h magnification Bright Field images of Fig. 4 show some atomic planes, but the resolution limit can not determine accurately the atomic plane thickness (Figure 7), the identified atomic planes belong to FeAl, Fe2Als and Fe AI3 c o m p a r e d to lattice parameters in Table 3.
Materials Processing and Texture
·
139
T E M S t u d y of t h e Inhibition L a y e r of C o m m e r c i a l H o t - D i p G a l v a n i z e d S t e e l s
Figure 7 : High magnification B F Images where we can see lattice parameters close to F e A l . FeAl and F e A b 2
5
Table III. S o m e Identified atomic planes from Figure 8 Lattice a, b, c (nm) <*, β, γ a=b=c= FeAl CC a = β = γ = 90° 0.289
Structure
a = 0.764 Fe AI 2
5
Orthorhombic
b = 0.641
a = β = γ = 90°
c = 0.421 a = 1.55 FeAlj
Monoclinic
b = 0.807
a = β = 90° γ = 120°
c = 1.247 DISCUSSION: T E M results confirm the double layer model reported in previous studies for samples galvanized at 0.16 and 0.18 w t % Al, prepared with chemical dissolution [7] and confirm also the presence of a Zn-alloy layer identified recently as F e Z n u (in addition to F e A l and F e A l i ) for samples galvanized at 0.20wt% and prepared by Focused Ion B e a m [17]. 2
140
Materials Processing and Texture
5
T E M S t u d y of t h e Inhibition Layer of C o m m e r c i a l H o t - D i p G a l v a n i z e d S t e e l s
A global schematic of the sample from steel substrate to the Zinc coating is shown in Figure 8. A b o v e the steel substrate, the first sublayer o f the inhibition layer is identified. It is composed o f a thin layer of crystallites with grain size around 4 n m (Figure 4-b). E D X analysis performed at different z o n e s on this surface and electron diffraction rings confirm that this layer is formed from a mixture of FeAl. Fe2Al and F e A I grains. This m a y be a result o f non-uniform dissolution of the Iron on the steel interface, or the Iron diffusion in FeiAl? grains. The presence of FeAl h a s not been identified before in the inhibition layer, but its presence is thermodynamically possible at 4 5 0 ° C according to the Ternary stable and metastable phase diagram of F e - A l - Z n [8]. A possible reason for not detecting FeAl layer In S E M studies is the xrays generation depth during the analysis (around 1 μηι) that is greater than the layer thinness (few n m ) . so the X-rays signals coming from the FeAl layer are mixed with the steel substrate signals. In the second sublayer, F e A l 5 grains were identified by diffraction rings (Figure 1 and 2), with s o m e traces of F e A l . Their diffraction rings are not continuous and less diffused than the first sublayer because o f big grain sizes, where most o f the spots belongs to the planes {221,002,101.311,130} of F e A l according to A S T M tables. T h e Fe2AI and F e A l j grains are surrounded by a Zinc-alloy layer, located between the inhibition layer and the Zinc coating, and having a strong Zn-Κα signal (Figure 3-d), this layer is a result o f reaction b e t w e e n the Zinc and the remaining dissolved Iron a t o m s that have not reacted with the Aluminum to form Fe-Al intermetallics. The Z i n c coating is not analyzed in this paper, it has usually the Z n - L greater than Zn-Κα signal in E D X analysis (contrary to the Zn-alloy layer), and the lower electrochemical potential (so, the priority in the electrochemical dissolution). It can be found only in some holes of the sample, but not on big surface. 5
3
2
3
2
5
5
a
Figure 8: Schematic of the inhibition layer morphology o f 0.35%-Al galvanized steel sheet CONCLUSION: In this paper, w e have studied galvanized steel sheet samples prepared by selective electrochemical dissolution. The samples were characterized by T E M imaging, diffraction and E D X on plane view. The double layer model o f the inhibition layer has been confirmed where w e c a n find the Fe-Al intermetallics arranged in t w o sublayers. F r o m the steel side, we can see the first sublayer containing small crystallites of 4 n m size and the roots o f F e A l and F e A l grains constituting the second sublayer. On the Zinc side, a Zn-alloy layer (Fe-Zn) with strong Zn-K,, signal surrounding the F e A l and F e A l grains has been identified and confirmed from the Zinc-coating traces on the surface. A schematic is finally d r a w n to present the inhibition layer morphology according to the previous pictures. 2
2
5
5
3
3
Materials Processing and Texture
141
T E M S t u d y of t h e Inhibition L a y e r of C o m m e r c i a l H o t - D i p G a l v a n i z e d S t e e l s
REFERENCES: 'A.R. M a r d e r - Progress in Materials Science 45 (2000) 191-271 Waldemar F U R D A N O W I C Z and C. R a m a d e v a S H A S T R Y - Metallurgical and Materials Transactions A, V O L U M E 30A, D e c e m b e r (1999) — PP. 3031-3044 J . S T R U T Z E N B E R G E R and J. F A D E R L - Metallurgical and Materials Transactions A, V O L U M E 29A, February 1998 — pp. 631-646 "Rodolphe P A R I S O T , Samuel F O R E S T , André P I N E A U , François G R I L L O N , Xavier D E M O N E T , and Jean-Michel M A T A I G N E - Metallurgical and Materials Transaction A, V 3 5 A . March 2004, P.813-823 ' R o d o l p h e P A R I S O T , Samuel F O R E S T , André P I N E A U , Franck N G U Y E N , Xavier D E M O N E T , and Jean-Michel M A T A I G N E - Metallurgical and Materials Transaction A, V 3 5 A , March 2004, P.797-811 R . Parisot, S. Forest, A-F. G o u r g u e s , A. Pineau, D. Mareuse - Computational Materials Science Vol. 19 (2000) p p . 189-204 Eric B A R I L and Gilles L ' E S P E R A N C E - Metallurgical and Materials Transaction A, V 3 0 A , March 1999, P.681-695 P . Perrot, J-C. Tissier, and J-Y. Dauphin. Zeitschrift fur Metallkunde, 8 3 , (11) (1992), pp. 786790. Jinichiro N a k a n o , Dmitri V. Malakhov, Shu Y a m a g u c h i , Gary R. Purdy - C o m p u t e r Coupling of Phase Diagrams and Thermochemistry 31 (2007) 1 2 5 - 1 4 0 Ν - Υ Tang - Journal of Phase Equilibria Vol. 21 N o . 1 (2000) " J . R . M C D E R M I D , M.H. K A Y E , and W.T. T H O M P S O N - Metallurgical and Materials Transaction Β Vol. 3 8 B , A P R I L 2 0 0 7 — pp.215-230 N - Y Tang - Metallurgical and Materials Transaction A, V 2 6 A , July 1995, P.1699-1704 P a t r i c k T O U S S A I N T . Luc S E G E R S , René W I N A N D and Michel D U B O I S - ISU International, Vol. 3 8 . N o . 9 , pp. 985-990 M - L . G I O R G I , R. N I C O L L E and J-B. G U I L L O T - Calphad, V o l . 2 5 , N o 3, pp 461-474, 2001 M - L . GIORGI and J - B . G U I L L O T - Journal of Materials Science 39 (2004) 5803 - 5808 M - L . G I O R G I , P. D U R I G H E L L O , R. N I C O L L E and J-B. G U I L L O T - Journal of Materials Science 40 (2005) 2263 - 2 2 6 8 3
6
7
8
9
ι α
l 2
l3
l 4
15
l 6
l7
E r i k a Bellhouse and Joseph M c D e r m i d - Galvatech 07 International Conference. 'Transmission Electron Microscopy V o l u m e II - Diffraction, C h a p 16: A Textbook Materials Science - by David B. Williams and C. Barry Carter. 1
142
Materials Processing a n d Texture
for
M E A S U R E M E N T O F T H E T E X T U R E S H A R P N E S S IN G R A I N - O R I E N T E D E L E C T R I C A L STEELS M. Frommert, C. Zobrist, D. Raabe, S. Zaefferer Max-Planck-Institut für Eisenforschung Max-Planck-Str. 1 40237 Düsseldorf, G e r m a n y L. Lahn, Α. Böttcher T h y s s e n K r u p p Electrical Steel Kurt-Schumacher-Str. 95 45881 Gelsenkirchen, G e r m a n y
ABSTRACT The measurement of the final Goss texture sharpness in grain-oriented electrical steels is a challenging task due to the immense grain size ranging from millimeters to centimeters. Although it is widely claimed in the literature that the orientation deviations from the ideal Goss orientation lie in the range of about seven degrees for conventional grain-oriented steel and about three degrees for high permeability grades, no precise investigation with an appropriate statistical relevance is k n o w n to the authors. In this work, X-ray diffraction ( X R D ) and large-area electron backscatter diffraction ( E B S D ) based orientation microscopy were used for texture analysis and orientation determination in order to estimate the Goss orientation spread of different grades of grainoriented steel. Additionally, t w o production routes for grain-oriented steel sheets are compared, the conventional route and a low heating route with lower inhibitor strength. The results of the texture measurement demonstrate that both routes deliver comparable values of orientation deviations. Furthermore, it can be shown that small differences in magnetic properties can be correlated with the texture sharpness of the material. INTRODUCTION Grain-oriented Fe-3%Si steels are used as cores in electrical transformers due to their specific soft-magnetic properties. They are characterized by a sharp {011} <100> texture, the socalled Goss texture, and a very large grain size in the range of millimeters to centimeters which develops at the end of a complex production process during discontinuous grain growth. The texture sharpness is usually regarded as a quality measure for the magnetic properties of grain-oriented steels (and vice versa) and misorientation values of 3 and 7 of the <100>axis from the rolling direction are generally accepted for high permeability grain oriented (HGO) and conventional grain oriented ( C G O ) grades, respectively " . The analysis of texture sharpness is not at all trivial due to the e n o r m o u s grain size and has been a matter of neglect in recent years although m o r e powerful tools with regard to measurement speed and precision have been developed. Concerning the material, increasing quality d e m a n d s as well as the pressure to reduce manufacturing costs are the main driving forces for industrial research and development activities. N e w production techniques have been recently developed, such as low heating routes and thin slab casting m e t h o d s . The magnetic properties of these materials are of major interest 0
1
0
3
4
143
M e a s u r e m e n t of t h e T e x t u r e S h a r p n e s s in G r a i n - O r i e n t e d Electrical S t e e l s
to the steel producers and can be measured with standardized m e t h o d s , but little is k n o w n about the texture sharpness of the final sheet produced by either of these process routes. T h e present paper aims at presenting t w o methods for the measurement of texture in large grained materials based on X-ray diffraction ( X R D ) and electron backscatter diffraction ( E B S D ) in the scanning electron microscope ( S E M ) . In the authors v i e w the X R D method has the potential to become a standard method for quality control of Goss-textured electrical steel sheets.
Fig. 1. S a m p l e setup for X R D measurement: a stack of sheets is m o u n t e d in a holder for sample preparation and texture m e a s u r e m e n t o n the cross section ( R D : rolling direction, T D : transverse direction, N D : normal direction). EXPERIMENTAL The material used in this investigation w a s final sheet of Fe-3%Si H G O and C G O steel grades. Furthermore, t w o production routes were analyzed, the conventional route (route 1) and a low heating route (route 2) with lower inhibitor strength. All material was kindly provided by ThyssenKrupp Electrical Steel. In order to check the accuracy of the X R D setup a Goss-oriented Fe-3%Si single crystal, produced at the Max-Planck-Institut für Eisenforschung by zone melting, was also investigated. X R D measurements were carried out in a goniometer equipped with a sample stage with 6 degrees of freedom, namely an Eulerian cradle with φ-, χ- and co-rotation in combination with a translation sample holder with x-. y- and z-movement. At every pole figure measurement position the sample was moved in the x- and y-direction in order to cover the m a x i m u m possible area of the measured sample. For detection an area detector with a high measurement speed (200 counts/pixel-s) w a s used. The setup further consisted of a cobalt X-ray tube and a polycapillarycollimator for the generation of a parallel b e a m with 1 m m spot size. For covering large sample areas, the final sheet w a s cut by spark erosion into one hundred slices of 10x30 m m , stacked and m o u n t e d in cross section in a holder for sample preparation and texture analysis (Fig. 1 ). With this method, probing of grains from m o r e than 4 0 0 c m was realized, containing about 1400 grains in the case of C G O (route 1) and about 130 grains in the case of H G O (route 1), respectively, thus providing excellent statistical data. 2
2
0
X R D analysis w a s performed using a standard grid of 5 for the φ-rotation. measuring the (100) and (110) pole figures which takes a total time of less than 4 hours. For determination of the texture sharpness, the peaks of the pole figures w e r e fitted along the φ-circle at β = 45 °
144
Materials Processing and Texture
M e a s u r e m e n t of t h e T e x t u r e S h a r p n e s s in G r a i n - O r i e n t e d Electrical S t e e l s
(β: pole distance) by a Gaussian distribution as s h o w n in Fig. 2. The value σ of the full width at half m a x i m u m ( F W H M ) represents the directional spread about the normal direction and is a measure for the orientation spread of the G o s s orientation.
0
Fig. 2. E x a m p l e of a 110-peak along the φ-circle at β = 45 fitted by a Gaussian distribution used for the determination of the Goss orientation spread by calculation of the full width at half maximum (FWHM). In contrast to the calculated orientation distribution function ( O D F ) obtained by X R D , the E B S D technique directly measures the O D F by means of individual crystal orientation measurements, therefore avoiding possible calculation errors. The use of a large scan E B S D setup enabled accurate analysis and rapid acquisition of the diffraction patterns. For the present investigation a J E O L 840A S E M working at an acceleration voltage of 20 kV was utilized. Large area scans on the sheet surface were realized in a c o m b i n e d stage and b e a m scanning m o d e where connected areas of 1700x600 μηι were approached via stage m o v e m e n t and the orientations then measured via beam m o v e m e n t with a step size of 100 μηι. The E B S D system was carefully calibrated to account for the pattern centre change occurring during the large area scans. The m e a s u r e m e n t took about 30 minutes for a single sample of 2 x 3 m m size, providing a total area of about 20 c m for the display of the pole figures for each material. It should be mentioned that this area contains only about 70 grains for the C G O (route 1) and about 10 grains for the H G O (route 1). 2
2
2
C o m p a r e d to the X R D measurement, the sample preparation for E B S D analysis is significantly more laborious and time-consuming since a perfectly polished surface is a prerequisite for a sufficiently good E B S D pattern quality. Particularly, the preparation of large sample areas o n the sheet normal as required for the investigation described above is a major difficulty of this technique. RESULTS AND DISCUSSION X R D measurements The pole figures plotted in Fig. 3 shovv the (110) poles of conventionally processed (route 1 ) H G O (a) and C G O (b) as well as a Goss oriented single crystal (c). A s expected, the
Materials Processing and Texture
·
145
M e a s u r e m e n t of t h e T e x t u r e S h a r p n e s s in G r a i n - O r i e n t e d Electrical S t e e l s
peak width, i.e. the orientation spread, is distinctly larger in the C G O material than in the H G O . The calculation of the F W H M - v a l u e s of the poles in the pole figures as described in Fig. 2 deliver orientation spreads of 9 ° and 6 ° for C G O and H G O . respectively. These results were obtained with a standard measurement grid of Δ φ =5 °. Furthermore, orientation deviations of 11 ° and 7 ° were obtained for C G O and H G O , respectively, processed via route 2.
Fig. 3. X R D pole figures of (a) H G O and (b) C G O (both route 1) measured with a grid of Δ φ = 5 ° and (c) pole figure of a Goss oriented single crystal measured with Δ φ = 1 °. All X R D analysis results, as summarized in Table 1, show a correlation with the magnetic properties, even for only slight differences in core loss and permeability between materials of the same grade processed via different routes. Table I. Comparison of the texture sharpness measured with X R D and E B S D .
The accuracy of the X R D setup w a s examined by an additional investigation of a single crystal with exact G o s s orientation. For the s a m e measurement grid ( Δ φ =5 °). an orientation spread of 5 ° was found which matches the magnitude of the grid. A change to higher orientation resolution, i.e. to a smaller grid of Δ φ =1 °. led to considerably smaller orientation deviations: in case of the single crystal to a F W H M - v a l u e of 1.1 ° (Fig. 3c). A high resolution measurement with the narrow grid was accomplished for the H G O material produced via route 1. and a value of 3.8 ° w a s found for the orientation spread w h i c h is markedly smaller than the measurement with the larger grid of Δ φ =5 °. as s h o w n above. Consequently, for the exact analysis of small orientation deviations a high orientation resolution, i.e. a fine grid, is necessary, w h i c h extends the measurement time by a factor of 3 to approximately 11 hours. H o w e v e r , the total time can be drastically reduced (by a factor of 4 or more) if only one peak of the pole figure is measured
146
·
Materials Processing a n d Texture
M e a s u r e m e n t of t h e T e x t u r e S h a r p n e s s in G r a i n - O r i e n t e d Electrical S t e e l s
w h i c h is sufficient for the analysis proposed above. M e a s u r e m e n t s following this scheme are currently performed in order to complete the X R D analysis with a fine grid and to fill out the voids in table I. Besides the predominant Goss orientation as the main texture component, in some pole figures a w e a k intensity s h o w s up which can be described as a fibre texture formed by a <100>rotation parallel to the rolling direction, therefore, belonging to texture components on the η fibre. In Fig. 4 two pole figures are plotted for H G O and C G O where these η-fibre intensities can be clearly seen.
Fig. 4. X R D pole figures of (a) H G O and (b) C G O (both route 1 ) showing weak intensities (green iso-intensity lines) on the η-fibre (<100> parallel to R D ) . 5
For the C G O material Chen et a l . investigated the effects of topology on abnormal grain growth. They could s h o w that the texture of the subsurface layer in the primary recrystallized microstructure is dominated by a weak γ-fibre, a texture c o m p o n e n t on the η-fibre ( 1 5 ° rotated from the cube orientation) and a minor 45 °-rotated cube orientation. However, the largest grains (or those with the highest n u m b e r of nearest neighbours) in the subsurface layer exclusively had orientations close to the η-fibre, the m a x i m u m orientation density being about 12 ° rotated from the G o s s orientation. These results suggest that the η-fibre found in the texture of the final sheet may be inherited through the secondary recrystallization process and, therefore, is a remnant of the primary recrystallized structure. Moreover, it is possible that some of those large grains can grow abnormally due to their size advantage. Further investigations into this aspect are extremely difficult d u e to the rareness of nucleation of secondary recrystallization. EBSD measurements A n example of E B S D m e a s u r e m e n t s obtained from single scans of H G O and C G O samples processed via route 1 is shown in Fig. 5. It is obvious that the grain size of H G O is significantly larger than that of C G O . Furthermore, the orientation deviation from the ideal Goss orientation is smaller in the H G O as can be seen from the coloring of the maps.
Materials Processing and Texture
·
147
M e a s u r e m e n t of t h e T e x t u r e S h a r p n e s s in G r a i n - O r i e n t e d Electrical S t e e l s
Fig. 5. E B S D images of the normal plane of H G O and C G O (both route 1); the color code indicates the deviation from the ideal G o s s orientation. In order to account for inaccuracies due to sample positioning and the reduced statistical significance of the data w h e n compared to the X R D method, the texture sharpness was not directly calculated as the deviation from the theoretical Goss orientation. Instead, a reference orientation well out of the Goss orientation spread w a s selected and the deviation from this theoretical orientation measured and classified into intervals of 1 °. From these data the cumulative frequency was calculated as s h o w n in Fig. 6a.
Fig. 6. Schematic diagram of the cumulative frequency of the deviation from the ideal G o s s orientation and resulting Gauss distribution used for the evaluation of the E B S D data.
148
·
Materials Processing a n d Texture
M e a s u r e m e n t of t h e T e x t u r e S h a r p n e s s in G r a i n - O r i e n t e d Electrical S t e e l s
Subsequently, an orientation distribution was computed from the cumulative curves under the following assumptions: 1. The true major orientation is the Goss orientation. 2. T h e deviation from the ideal G o s s orientation follows a Gaussian distribution. T h e appropriate distribution is s h o w n in Fig. 6b. The F W H M - v a l u e s determined from these curves are collected in Table 1. T h e texture sharpness tends to have smaller values - about 3 ° for H G O and 6 ° for C G O - as compared to the X R D results measured with Δ φ =5 °, but the same tendency is found concerning the difference b e t w e e n the t w o grades of electrical steel. Moreover, the result of the high resolution X R D analysis ( Δ φ =1 °) performed with H G O (route 1 ) is in good agreement with the result of the E B S D measurement. For better visualization, Fig. 7 displays the (110) pole figures for both grades of material processed via routes 1 and 2. The texture data w e r e obtained from the E B S D measurements after merging all maps belonging to one sample. A comparison of the pole figures shows larger orientations spreads and lower intensities for C G O than for H G O , but similar peak width for the same material processed via routes 1 and 2.
Fig. 7. E B S D pole figures of H G O and C G O for routes 1 and 2. CONCLUSIONS The Goss orientation spreads for H G O and C G O lie in the range of 3 ° to 6 °, respectively, w h e n measured by E B S D , and 6 ° to 11 ° when measured by X R D using a grid of Δ φ = 5 °. The discrepancy between the results of both measurement techniques, E B S D and X R D , could be attributed to the orientation resolution of the X R D grid: an orientation spread of only 3.8 ° w a s found for H G O using a grid of 1 ° which closely corresponds to the E B S D result. Besides, it could be s h o w n that both production routes - the conventional and the low heating process - deliver material of c o m p a r a b l e texture sharpness.
Materials Processing and Texture
149
M e a s u r e m e n t of t h e T e x t u r e S h a r p n e s s in G r a i n - O r i e n t e d Electrical S t e e l s
A weak η-fibre was found in the pole figures obtained by X-ray diffraction. This component may be attributed to inheritance m e c h a n i s m s from the primary recrystallized structure where the largest grains have orientations close to the η-fibre. T h e setup proposed for X-ray measurement offers an efficient tool for the determination of the texture sharpness of large grained materials due to excellent statistics and less experimental effort c o m p a r e d to E B S D measurements. The measurement time for the X R D analysis can be drastically reduced if only one peak of the pole figure is measured which is sufficient for the determination of the texture sharpness. Further experiments are currently being performed to complete the X R D measurements with high orientation resolution (Δφ = 1 °) in order to further correlate the texture sharpness with the magnetic properties of the materials processed by the different production routes. REFERENCES 'R. Ricci Bitti and M. Candiotti, "Surface Energy and Primary Texture Influence on the Improvement of Goss Texture in High-Permeability Electrical Steel", Journal of Magnetism and Magnetic Materials, 26, 11 -14 (1982). S . Mishra, C. Därmann, Κ. Lücke, "On the Development of the Goss Texture in I r o n - 3 % Silicon", Acta Metall., 12, 2185-2201 (1984). M . Matsuo, "Texture Control in the Production of Grain Oriented Silicon Steel", ISU International, 2 9 , 809-827 (1989). K . Günther, G. A b b r u z z e s e , S. Fortunati, G. Ligi, "Recent Technology D e v e l o p m e n t s in the Production of Grain Oriented Electrical Steel", Steel Research Int., 7 6 , 4 1 3 - 4 2 1 (2005). N . Chen, S. Zaefferer, L. Lahn, Κ. Günther, D. Raabe, "Effects of topology on abnormal grain growth in silicon steel", Acta Mater., 5 1 , 1755-1765 (2003). 2
3
4
5
150
Materials Processing a n d Texture
EFFECTS OF COMPOSITION A N D COILING TEMPERATURE ON PRECIPITATION AND T E X T U R E F O R M A T I O N IN A F E W I N T E R S T I T I A L F R E E H I G H S T R E N G T H S T E E L S Pampa Ghosh and R Κ Ray R & D Division, Tata Steel Jamshedpur, Jharkhand, India ABSTRACT Interstitial free high strength (IFHS) steels are widely used in the automobile sector for their excellent deep drawability coupled with adequate strength. However, very often deterioration of drawability has been reported in these steels, as compared to the normal interstitial free (IF) steels. A few batch annealed I F H S steels with varying chemical composition and coiled at different temperatures have been chosen for the present study with a view to determining the effects of composition and coiling temperature on precipitation and texture formation after cold rolling and annealing. It has been observed that the final texture developed after batch annealing depends on precipitation which in turn can be related to the composition and processing history of the steels. In these steels Ρ and M n were added as solid solution strengthening elements, w h e r e a s Ti and N b additions were m a d e to scavenge the C and Ν in solid solution. It appears that Ρ is taken out of the solution due to the formation of FeTiP, thus leaving behind insufficient Ti to form T i C . As a result C remains in solid solution thereby degrading the beneficial {111} texture formation during recrystallization annealing. However, w h e n the s a m e steel is coiled at a higher coiling temperature, precipitation of carbosulphides takes place during the coiling itself which in turn takes the C out of solid solution leading to an enhancement of the favorable {111 (texture formation during recrystallization anneal. Presence of FeTiP and absence of carbide type of precipitates have been identified as the key reasons for the deterioration of deep drawability in these steels w h e n coiled at lower temperatures. INTRODUCTION IFHS-steels, like IF steels are normally processed by the conventional Hot Rolling, Cold Rolling and Annealing route. Annealing after cold rolling can be done by either the batch or continuous process. Every processing step in the production of these steels, from composition control to hotmill processing, cold rolling and annealing, has been found to affect the properties of these steels . There is every reason to believe that the texture formation and mechanical behavior of such steels could be very m u c h related to their precipitation behavior. This present study has been undertaken with an aim to find out the effect of composition and coiling temperature of hot rolling on the precipitation behavior and texture formation in a few interstitial free high strength steels. Attempt has also been m a d e to establish correlation (if any) between precipitation and texture formation in the same. 1
EXPERIMENTAL A total of four industrially batch annealed I F H S steels were chosen for this study. Out of these four, Steels 1 and 2 were from the same cast having the same composition. However, their coiling temperatures were different. Steels 3 and 4 are from different casts and are of different compositions, while both of them were coiled at the same temperature. The chemical compositions and processing parameters of the four steels are given in Tables I and II, respectively.
151
Effects of C o m p o s i t i o n a n d Coiling T e m p e r a t u r e o n P r e c i p i t a t i o n a n d T e x t u r e F o r m a t i o n
Table I: Chemical Compositions of Experimental Steels (in w t % , * in ppm) Steel 1 2 3 4
C* 40 40 38 40
Mn 0.84 0.83 0.75 0.50
Si 0.018 0.020 0.020 0.014
S* 70 70 70 90
Ρ 0.091 0.090 0.086 0.058
AI 0.048 0.049 0.050 0.029
Ti 0.070 0.069 0.065 0.041
Nb 0.030 0.030 0.029 0.018
N* 38 38 40 38
Table II: Processing Parameters of Experimental Steels Cold Rolled Batch Annealing Furnace Temperature (°C) 700 700 700 700
Hot Rolled
Steel
1 2 3 4
Finish Rolling Temperature ( F R T , °C)
Coiling Temperature (CT, °C)
Cold Roll Reduction (%)
910 910 910 910
550 620 550 550
54 52 54 72
Skin Pass Mill Elongation (%) 0.5 0.5 1.2 1.2
The as received steel samples were examined in an optical microscope and in a Scanning Electron Microscope ( S E M ) equipped with an Energy Dispersive Spectrometer (EDS) to reveal the overall structures and bigger size precipitates (>0.5 μπι). For the characterization of n a n o sized precipitates, a Transmission Electron Microscope ( T E M ) equipped with E D S was used. Specimens for T E M were prepared by carbon extraction replica as well as by the thin foil method to study the morphology, size distribution and chemistry of the precipitates. Crystallographic texture w a s determined from the mid thickness regions of the as received samples using a Panalytical half circle texture goniometer and (110), (200), (112) and (222) pole figures were measured. Orientation distribution functions ( O D F ) were then calculated from these incomplete pole figures using the M T M - F H M software . T h e mechanical properties such as yield strength, ultimate tensile strength and percentage elongation, along with average r value, were measured and the data are listed in Table III. 2
Table III: Mechanical Properties of the Experimental Steels Steel
YS ( M P a )
U T S (MPa)
%E1
1 2 3 4
278 258 320 218
461 419 448 391
35 42 36 39
r 1.28 1.90 1.37 1.62
RESULTS All the four steels s h o w well recrystallized grains in the final annealed condition as shown in Figure 1. The measured average grain sizes of the steels (Steels 1, 2, 3 and 4) were 5.30, 8.77, 6.08 and 5.34 μπι, respectively. Steel 2 shows perceptibly larger grain size as
152
Materials Processing and Texture
Effects of C o m p o s i t i o n a n d Coiling T e m p e r a t u r e o n Precipitation a n d T e x t u r e F o r m a t i o n
compared to the other three steels and the reason is not hard to understand, as this steel was coiled at a higher temperature as c o m p a r e d to the rest.
(c) (d) Figure 1 : Optical micrographs of cold rolled and annealed steels; (a) Steel 1, (b) Steel 2, (c) Steel 3 and (d) Steel 4. S E M study of all the steels reveal the presence of rather large cuboidal precipitates o f square or rectangular cross-section. E D S study identified these precipitates as TiN. In most cases, these precipitates are found to be present in association with A 1 0 particles. It may be mentioned that these steels are Al-killed a n d Ti and N b were added once the steels were killed. As a result formation o f TiN particles (which may form w h e n the steel is still liquid) may occur heterogeneously over existing AI2O3 particles acting as preferential nucleation sites. Occasionally M n S precipitates are observed to be present on TiN particles. The formation temperature of M n S is much lower a s c o m p a r e d to that of TiN and it always forms in the solid state. H o w e v e r , free standing M n S has also been observed. There is not m u c h difference in the precipitation behavior of the four experimental steels w h e n observed under the S E M and some typical features are shown in Figure 2. 2
3
Materials Processing a n d Texture
·
153
Effects of C o m p o s i t i o n a n d Coiling T e m p e r a t u r e o n Precipitation a n d T e x t u r e F o r m a t i o n
(e) (0 (g) (h) Figure 2: S E M m i c r o g r a p h s of the steels showing presence of bigger sized precipitates; (a-c) Steel 1. (d-e) Steel 2, (f-g) Steel 3 and ( 0 Steel 4. T E M observations, however, h a v e revealed significant differences in the precipitation behavior of the four steels. Very fine precipitates, their sizes ranging between 10-500 nm, are found to be present in all the steels, although their nature and v o l u m e fractions differ from steel to steel. The highest v o l u m e fraction of precipitates has been observed in Steel 1 followed by Steels 2, 3 and 4 in that order. A l t h o u g h no conscious effort w a s m a d e to m e a s u r e the v o l u m e fraction o f precipitates quantitatively, the present results relate to observations of a very large n u m b e r of extraction replicas taken from each steel. The overall features of the fine precipitates in the various steels are s h o w n typically in Figure 3. Very often clusters of precipitates are observed as shown in Fig. 3(b). This clustering is most c o m m o n in Steel 1 as c o m p a r e d to the rest. Careful observation of the micrographs also reveal the presence of very fine precipitates (<10 nm) in Steel 1. Using E D S and by analyzing the diffraction patterns most of the precipitates in these steels were identified as F e ( T i + N b ) P . These precipitates do not have any definite shape and c a n be square, lenticular or globular. Very often these precipitates are observed to be present epitaxially on carbide or sulphide precipitates. Thin foil T E M study o f the steels has also s h o w n the presence of these precipitates on grain boundaries thereby causing grain boundary tearing. All these features are shown in Figure 4. Comparison of the precipitation behavior of Steels 1 and 2 reveals that although most of the precipitates in both the steels are o f F e ( T i + N b ) P type, Steel 2 occasionally s h o w s the presence o f T14C2S2 (Figure 4(c-d), w h e r e a s Steel 1 d o e s not. N o n e of these t w o steels show any evidence of carbide precipitation. C o m p a r i s o n of the precipitation behavior of Steels 3 and 4 reveals the presence of T i C and F e ( T i + N b ) P type o f precipitates in both [Fig 4(a-b)]. It appears that the v o l u m e fraction of T i C precipitates is higher in Steel 4 as c o m p a r e d to Steel 3. A s c o m p a r e d to Steel 1, the total v o l u m e fraction o f F e ( T i + N b ) P type o f precipitates appears to be less in Steel 3 and very much less in Steel 4. The overall precipitation behavior of all the four steels is summarized in Table IV.
154
Materials Processing a n d Texture
Effects of C o m p o s i t i o n a n d Coiling T e m p e r a t u r e o n P r e c i p i t a t i o n a n d T e x t u r e F o r m a t i o n
Figure 3: T E M micrographs of the experimental steels showing overall features of precipitates: (a-b) Steel 1, (c-d) Steel 2 . (e) Steel 3 and (f) Steel 4.
Materials Processing a n d Texture
155
Effects of C o m p o s i t i o n a n d Coiling T e m p e r a t u r e o n Precipitation a n d T e x t u r e F o r m a t i o n
Figure 4 : T E M micrographs o f the experimental steels s h o w i n g different precipitates in (a) Steel 3. (b) Steel 4 (c) Steel 2, (d) S A D P o f T14C2S2 in Steel 2, (e) grain boundary precipitate in Steel 1 and (f) grain boundary precipitate in steel 4.
156
Materials Processing a n d Texture
Effects of C o m p o s i t i o n a n d Coiling T e m p e r a t u r e o n P r e c i p i t a t i o n a n d T e x t u r e F o r m a t i o n
Table IV: Overall Precipitation Behavior of the Experimental Steels Precipitates
Steel 1
Steel 2
Steel 3
Steel 4
Total V o l u m e fraction of precipitates Fe(Ti+Nb)P
Highest
Less than Steel 1, but higher than Steel 3 Present, most of the precipitates are of this type
Less than Steel 2 but higher than Steel 4 Present, most of the precipitates are of this type
Lowest
Present, most of the precipitates are of this type
Ti C S occasionally found Present Present
TiC occasionally found
TiC very often found
Present Present
Present present
Present, most of the precipitates are of this type. Very fineness than 10 n m ) ppt. ofFe(Ti+Nb)P are very often observed throughout the matrix
Carbide/carbosulphide
Not found
TiN MnS
Present Present
4
2
2
Figures 5 (a-d) show the Φ = 45° sections of the O D F s of the steels. It is clear that Steel 2 has the strongest {111} texture c o m p o n e n t followed by Steel 4, Steel 3 and Steel 1, in that order The measured average r values also follow the same pattern. Figure 6 s h o w s the γ-fiber intensity plots for all the four steels and invariably Steel 2 shows the highest intensity followed by Steels 4, 3 and 1. 2
3.20 1.30
2.50 1.00
2.00 1.60 0.80
8.00 2.00
5.60 1.40
4.00 1.00
2.80 0.70
8.00 3.20 1.30
6.40 2.50 1.00
5.00 4.00 2.00 1.60
Π . 0 8.00 2.80 2.00 0.70
5.60 1.40
4.00 1.00
(a) (b) (C) (d) Figure 5: Φ = 45° sections of the O D F s of the steels; (a) Steel 1. ( b ) Steel 2. (c) Steel 3 and (d) Steel 4. 2
Materials Processing and Texture
157
Effects of C o m p o s i t i o n a n d Coiling T e m p e r a t u r e o n Precipitation a n d T e x t u r e F o r m a t i o n
Figure 6: γ-fiber intensities of the experimental steels
DISCUSSION The precipitation behavior and texture of the experimental steels clearly reveal that there is a strong correlation b e t w e e n precipitation and texture formation. T h e total volume fraction of precipitates in these steels appears to increase with increasing Ρ and Mn content. As the precipitates are mainly of the FeTiP type, a higher Ρ content is expected to increases the tendency of F e T i P formation. This will result in a depletion of Ti for TiC formation or at least this will be responsible for a competition between TiC and FeTiP formation ' . A s hardly any carbide precipitation has been observed in Steel 1, it could be due to the presence of higher amount of Ρ in that steel, leading to copious FeTiP formation, thereby hardly leaving any Ti to tie up with carbon. A s a result carbon remains in solid solution and degrades the {111J texture formation, thereby resulting in a poor average value of r . In case of Steels 3 and 4, lesser amount of Ρ causes a decrease in the a m o u n t of F e T i P precipitation, thereby allowing m o r e carbide precipitation. Steel 4 contains the lowest amount of Ρ a m o n g the three steels (Steels 1, 3 and 4), coiled at the lower temperature of 550°C. This also s h o w s the highest average r value a m o n g the three steels. Although, Steels 1 and 2, have similar chemical composition, they exhibit very different textures. This appears to be due to the difference in the coiling temperature of the steels. Steel 2 which was coiled at the higher temperature of 6 2 0 ° C s h o w s the presence of carbo-sulphide precipitates during coiling itself. T h u s carbon is taken out of solid solution before the onset of formation of FeTiP precipitates. This will result in a stronger {111} texture formation in Steel 2 leading to the achievement of the highest average r-value a m o n g the four steels. The results of this investigation clearly s h o w that both composition and coiling temperature of I F H S steels affect the sequence and nature of the precipitation behavior and textural evolution. 3
158
Materials Processing and Texture
4
Effects of C o m p o s i t i o n a n d Coiling T e m p e r a t u r e o n Precipitation a n d T e x t u r e F o r m a t i o n
CONCLUSIONS The following conclusions can be made out of this study: • Carbon in solid solution itself is very harmful for {111} texture formation. • FeTiP precipitation is harmful for {111} texture formation, specifically if it forms before TiC formation. FeTiP precipitation takes out Ti form the solid solution and makes the matrix lean of Ti to tie up with carbon. • Increasing Ρ content and keeping all the other factors constant, increases the volume fraction of FeTiP precipitation, thereby degrading the {111} texture. • Coiling at higher temperature, but below FeTiP formation temperature, improves the {111} texture formation due to the formation of carbo-sulphide type precipitates during the coiling stage.
REFERENCES ' S . Holie, Processing and Properties of Mild Interstitial Free Steels, Materials Science and Technology, 16, 1079-1093 (2000). P. Van Houtte, The M T M - F H M Software System, Version 2, Department of Metallurgy and Materials Engineering, Katholieke Universiteit Leuven, Belgium. O. Hamart, T. Lung, S. Lanteri, MWSP ConfProc, 40, 189 (1998). T. Gladman, Physical Metallurgy of Microalloyed Steels, Springer, Berlin (1977) P . Ghosh, R. K. Ray, B . Bhattacharya, S. Bhargava, Precipitation and Texture Formation in T w o Cold Rolled and Batch Annealed Interstitial Free High Strength Steels, Scripta Materialia, 55, 271-274 (2006) 2
3
4
5
Materials Processing and Texture
159
This Page Intentionally Left Blank
M I C R O S T R U C T U R E A N D T E X T U R E D E V E L O P M E N T IN W A R M R O L L E D HIGH S T R E N G T H IF S T E E L S C O N T A I N I N G Cu, Ρ A N D Β Arunansu Haldar and Chiradeep Ghosh R & D Division, Tata Steel Jamshedpur, Jharkhand, India ABSTRACT W a r m rolling or ferritic rolling of low carbon steels and interstitial free steels is gaining importance due to its potential to broaden the product range and reduce the production cost. The present paper discusses the evolution of microstructures and textures in Cu, Β and Ρ containing IF steels after multipass warm rolling with one finish rolling temperature (FRT) and fixed amount of deformation in ferritic region. T h e effect of warm rolling on Cu precipitation during ageing was also studied. It has been found that like cold rolling, warm rolling does not affect the kinetics of the C u precipitation to that extent because of the occurrence of Cu precipitation during rolling and cooling. Very strong γ and α textures were developed after warm rolling in both the steels. In order to further strengthen the γ texture and also to extract the maximum benefit of precipitation, both the steels were given special heat treatment. T E M study was carried out in order to characterize the microstructures and Cu precipitation. INTRODUCTION T h e rolling of steel strips at a temperature below A r temperature (800°C - 600°C) is known as warm rolling ( W R ) or ferritic rolling. This has been studied extensively nearly for a decade by several workers to understand the physical metallurgy and advantage of warm rolling particularly for IF steel " . W R provides considerable amount of savings due to requirement of low reheating temperatures, less deformation energy (ferrite is softer phase), and lower scale loss and roll w e a r . T h e W R was found to be very effective for IF steel in producing even formable sheet comparable to the conventional cold rolling and annealing route. But low carbon steel (LC) exhibits a very poor γ fibre texture " leading to low mean r-values with high directional variabilities (Ar) resulting poor formability . IF steel is well known for exhibiting very high formability (r-mean > 2.0, Ar ~ 0) but limits its application due to lower strength of this steel. Various strengthening alloying elements are added in this steel to increase the strength and applicability of this steel. Mn and Ρ are well known and effective solid solution strengthening elements are added t o IF steel t o increase its strength. T o counteract the cold embrittlement due to Ρ addition, boron is added to the steel. Cu is also another elements which is added to steels to provide further strengthening through C u precipitation by aging treatment. The present work describes the evolution of microstructures and textures of two IF steels: one containing only Cu and other containing Μ η , Ρ, Β and Cu during multi pass warm rolling. 3
1
4
2
5
7
8
EXPERIMENTAL PROCEDURE T w o IF steels containing Cu and or Μ η , Ρ, Β were taken for this study (Table 1). The steels were vacuum induction melted. The cast ingots were forged into the bar of 50 m m width and 12 m m thickness.
161
M i c r o s t r u c t u r e a n d T e x t u r e D e v e l o p m e n t in W a r m R o l l e d H i g h S t r e n g t h IF S t e e l s
Table I. Chemical composition of the steels (wt%) Ρ Si Mn S Steel C Steel 1
IF-Cu (AC|=880°C Ar =842°C) IFHS-Cu (Ac,=820°C Ar.,=640°C)
Al
B
0.003
0.12
0.01
0.016
0.50
0.023
Cu 1.04
Cr 0.025
Ti 0.063
0.008
N 0.0026
0.003
0.34
0.01
0.083
0.50
0.020
1.02
0.025
0.066
0.011
0.0023
3
Steel 2
These bars (150 m m X 5 0 m m X 12 m m ) were used as the feed stock for warm rolling. T h e bar was soaked at 1145 °C and then rolled as per the following schedule:
nJ
l"Pass
12 m m
2 Pass
8 mm 33%
3
•
r d
Pass
5 mm
37.5%
•
3 mm
40%
The first pass was finished within 1000 °C and within austenite region. Last t w o passes were given within 800 - 675 °C. T h e finish rolling temperatures were 7 0 0 ° C and 650°C for steel 1 and steel 2 respectively. Rolling temperatures were controlled in such a way that last two passes were given in ferritic region for IF-Cu steel and in two phase region ( γ + α ) for IF-HS Cu steel (considering their Arj temperatures determined at cooling rate of l°C/s as shown in Table I). The surface temperature of the bar was measured before and after each pass using laser optical pyrometer. After rolling the steel strip were water quenched immediately within 5 seconds to retain the deformation state of the strip. S o m e samples were given ageing treatment at the temperatures between 300 - 5 5 0 ° C for different times up to 300 minutes. Hardness values were measured for each sample. A two stage heat treatment was also performed with warm rolled samples for both the steels. For steel 1 the strip was heated to 800 °C for 10 minutes and immediately placed in a salt bath furnace at a temperature of 500 °C for 60 minutes and then water quenched. For the steel 2 the sample was heated to 725 °C for 20 minutes and followed the same process like steel I. Microstructures of the strips were studied on the longitudinal section containing normal direction (ND) and rolling direction (RD) using optical, Scanning electron microscope ( S E M ) and 120 KV Philips transmission electron microscope ( T E M ) . T E M sample was prepared in the conventional way by mechanical polishing (up to 0.1mm) and then by twin jet polisher using a mixture of glacial acetic acid and perchloric acid (80:20) solution as electrolyte. Electro polishing was carried out at 60-70V and at a temperature of about 12 °C. Texture at the mid thickness plane was measured for w a n n rolled samples and two heat treated samples using Philips M R D texture goniometer. Orientation Distribution Functions (ODFs) were calculated from (110), (200), (112) and (310) pole figures using the Method of B u n g e with the help of M T M software developed by Prof. Paul VanHoutte of K U , Leuven, Belgium. Tensile testing of the W R and heat treated samples were carried out in rolling direction of the samples at a cross head speed of 10 m m / m i n u t e using an Instron tensile testing equipment. 9
RESULTS AND DISCUSSION
162
Materials Processing and Texture
M i c r o s t r u c t u r e a n d T e x t u r e D e v e l o p m e n t in W a r m R o l l e d H i g h S t r e n g t h IF S t e e l s
T h e Figure 1(a) shows the microstructures of w a r m rolled sample of steel I. The microstructure consists of elongated ferrite grains which contain deformation bands aligned in specific direction. Figure 1(b) reveals these bands clearly when they are observed at higher magnification.
Fig.l showing (a) deformation bands formed within the grains and inclined at three different angles with respect to rolling direction as marked by arrows and (b) deformation bands inclined at one angle within one grain at higher magnification.
Materials Processing and Texture
163
M i c r o s t r u c t u r e a n d T e x t u r e D e v e l o p m e n t in W a r m R o l l e d H i g h S t r e n g t h IF S t e e l s
The microstructures as shown in Figure I is typical of IF steels when they are w a n n rolled. M a n y researchers also reported the s a m e occurrence of deformation bands and or micro bands in w a r m rolled condition ' ' . In fact, extensive T E M work by Haldar et. ai has shown that these extended planar boundaries starts forming at very early stage of deformation ( 1 0 % and 3 0 % deformation) and these planar boundaries are inclined at an angles of ± 35 to R D and they are all crystallographic in nature . T w o kinds of slip systems are operative at these temperatures and are decided by the higher Schmidt factor. Clear orientation effects o n microstructures w e r e also observed. But at higher rolling deformation (80%) at room temperature the lamellar structures form and they are almost parallel to rolling d i r e c t i o n . S o m e T E M work was also carried out on the longitudinal sections. But edge on condition could not be maintained and rolling direction of the thin foil w a s also not known on the sample holder. This is the reason w h y sharp boundaries could not be observed and their inclination angles with respect to rolling direction could not be measured. Figure 2 (a) and (b) show T E M micrographs of steel 1 at warm rolled condition. This shows deformation bands and large number of precipitates within the b a n d s along with large amount of dislocation densities. 1
, -
6
0
10
Fig. 2: S h o w i n g T E M micrographs of steel 1 from t w o different locations containing deformation bands and large number of precipitates.
164
Materials Processing and Texture
M i c r o s t r u c t u r e a n d T e x t u r e D e v e l o p m e n t in W a r m R o l l e d H i g h S t r e n g t h IF S t e e l s
Fig.3: T E M micrograph of steel I (a) at higher magnification showing Cu precipitates along with dislocations formed during warm rolling of steel 1 and (b) the corresponding E D S plot from the precipitates indicating those are Cu precipitates.
Materials Processing and Texture
165
M i c r o s t r u c t u r e a n d T e x t u r e D e v e l o p m e n t in W a r m R o l l e d H i g h S t r e n g t h IF S t e e l s
Figure 3 (a) shows Cu precipitates at higher magnification and E D S taken from this area shows the precipitates contain only Cu. Figures 4 (a) & (b) show the T E M micrographs of warm rolled microstructures of the steel 2. Similar structures as shown in 3 (a) and (b) were also evidenced in this steel (Fig. 4). Deformation bands with planar boundaries with lots of Cu- precipitates within are visible clearly in t w o micrographs taken from t w o different areas at two different magnifications as shown in Figures 4 ("a) and (b). Large n u m b e r of C u precipitates both finer and coarse are noticed, which were formed during warm rolling and cooling. This steel was finished rolled at lower temperature (650°C) as the A n temperature of this steel w a s lower (640°C).The precipitates were confirmed by E D S taken from this areas shown in the Figure 4 (c). Tn both the steels other kinds of precipitates of Ti were also observed which are not reported here.
Fig.4: T E M micrographs (a) and (b) of steel 2 taken from two different locations at different magnification showing deformation bands with planar boundaries and C u precipitates formed during warm rolling and (c) the corresponding E D S plot from the precipitates confirming those as Cu - precipitates.
166
Materials Processing and Texture
M i c r o s t r u c t u r e a n d T e x t u r e D e v e l o p m e n t in W a r m R o l l e d H i g h S t r e n g t h IF S t e e l s
Both the steels developed deformed microstructures with planar dislocation boundaries and cell structures. These are similar to the geometrically necessary boundaries ( G N B s ) and incidental . These dislocation boundaries (IDBs) as observed by others in F C C metals and steels ' microstructures are also similar to the deformation structures observed in IF steels by other w o r k e r s ' . Large n u m b e r of copper precipitates also formed during w a r m rolling, which were clearly evidenced in both the steels. T h e mechanical properties (average of three samples) of the warm rolled steels were measured in the rolling direction of the samples, which are shown in the Table II. 3
0,
11
L1
M
Table Π: S h o w s tensile properties of t w o steels in warm rolled conditions. Steel 1 2
YS, MPa 375.0 540.00
UTS,MPa 400.0 590.0
%E1 20 12
T h e mechanical properties of both the steels show the behavior similar or close to cold rolled steels exhibiting higher strength and lower elongation. Higher strength of both steels was contributed by warm rolled microstructures similar to cold rolled steels and Cu - precipitates formed during rolling and cooling. Deformed microstructures with large dislocation density were mainly responsible for the lower ductility. Because of the presence of large number of Cuprecipitates the ageing responses of both the steels were not very significant and consistent. Figures 5 (a) and (b) show the ageing curves (hardness vs. time) for both the steels. These do not give any correlation between hardness values, ageing temperatures and ageing time.
Fig.5: S h o w i n g hardness vs. time plots for two steels obtained after ageing treatments at different temperatures after warm rolling.
Materials Processing and Texture
·
167
M i c r o s t r u c t u r e a n d T e x t u r e D e v e l o p m e n t in W a r m R o l l e d H i g h S t r e n g t h IF S t e e l s
Texture development after w a n n rolling of both the steels was studied. The O D F plots of both the steels are s h o w n Figures 6 and 7.
Fig.6: O D F plot of φ , sections (0 - 90°) and (p = 4 5 ° section of warm rolled I F - C u steel indicating a strong texture development during w a n n rolling. 2
168
·
Materials Processing a n d Texture
Fig.7: O D F plot of φ , sections (0 - 90°) and (p = 4 5 ° section of warm rolled IFHS - Cu steel indicating a strong texture development during warm rolling. 2
M i c r o s t r u c t u r e a n d T e x t u r e D e v e l o p m e n t in W a r m R o l l e d H i g h S t r e n g t h IF S t e e l s
Figures 6 and 7 show that strong textures developed in both the steels as was observed by others " . Both α and γ fibre textures were found to be equally strong for both the steels which is 5
7
clearly evidenced in 92 = 4 5 ° section. Table III s h o w s the volume fractions of different texture components calculated using M T M - F H M software developed by Prof. Paul Van Houtte of KU, Leuven, Belgium. Table III: S h o w s calculated volume fraction of different texture c o m p o n e n t s (%) and r bar Steel
V o l u m e fraction of various texture c o m p o n e n t s , %
Calculated r bar
Gamma fibre
Alpha fibre
F
Ε
Η
I
Goss
{111}<112>
{111}<110>
{001}<110>
{112}<110>
{011}<100
1
25.0
33.0
12.0
12.0
5.0
15.5
3.0
1.30
2
23.5
33.5
11.5
11.0
6.0
16.0
3.0
1.10
Table III also shows that almost similar textures with minor variation developed in both the steels during warm rolling. Alpha fibre components are very strongly present in both the steels. be very strong. Presence of these α fibre T h e I component ( f l l 2 } < 1 1 0 > ) is found to c o m p o n e n t s and deformed structures are responsible for poor r bar values (calculated using Taylor Model) for both the steels. In order to improve the r bar values of the steels, a separate heat treatment w a s performed for both the W R steels. T h e warm rolled samples were first heated to ferritic region (below Aci temperature). T h e samples were first heated to the temperatures of 800 °C (steel 1) and 725 °C (steel 2) and kept at those temperatures for 15 minutes in order to develop strong γ fibre and also to get m o r e Cu into solution. After soaking the samples were immediately taken out and placed into a salt bath furnace maintaining a bath temperature at 550°C and aged at that temperature for o n e hour for Cu precipitation and then water quenched. The heat treated microstructures are shown in Figures 8 (a) and (b). T h e heat treated microstructures consist mainly of recrystallised ferrite grains. Polygonal ferrite grains are clearly visible in steel 1. O n the contrary, in steel 2 the microstructures were very fine [Fig. 8 (b)] and fully recrystallised ferrite. Hardness values for both the steels improved significantly after heat treatment. T h e hardness values IF-Cu steel went up to as high as V P N 4 7 0 from V P N 250 and from V P N 238 to V P N 4 9 0 for I F H S - C u steel, which are m o r e than what were observed after ageing the warm rolled samples. This clearly indicates that large amount of Cu precipitation took place during or after warm rolling and some of those got dissolved during first stage of heat treatment and then again re-precipitated during second stage of heat treatment at 550°C leading to an increment of hardness value.
Materials Processing and Texture
·
169
M i c r o s t r u c t u r e a n d T e x t u r e D e v e l o p m e n t in W a r m R o l l e d H i g h S t r e n g t h IF S t e e l s
Fig. 8: Shows optical microstructures of (a) IF-Cu and (b) 1FHS-Cu steels after second heat treatment After heat treatment similar kinds of textures containing both strong α and γ fibres were observed in both the steels. T h e calculated r-bar values neither deteriorated nor improved significantly. T h e calculated r-bar values were found to be 1.40 and 1.20 for IF-Cu and 1FHS-Cu steel respectively. In the conventional cold rolling and annealing process, the r-bar value of the IF-Cu steel could be improved to as high as 2 . 0 by heat treatment. Similarly R a n a et. a/.."'" also could improve the r-bar of 1FHS-Cu steel up to 1.77 in batch annealed condition, but in other conditions similar r-values, as obtained here, were reported. L
170
Materials Processing and Texture
18
M i c r o s t r u c t u r e a n d T e x t u r e D e v e l o p m e n t in W a r m R o l l e d H i g h S t r e n g t h IF S t e e l s
Fig.10: O D F plot of (pisections (0 - 90°) and φ = 4 5 ° section of warm rolled and heat treated (a) IF-Cu and (b) IFHS - Cu steels indicating a strong texture development after heat treatment. 2
CONCLUSIONS T w o kinds of IF steels containing Cu were found to develop moderately sharp texture having equal amount of α and γ fibre components in warm rolled condition. Strength of the steels could be improved due to warm rolling because of the presence of deformation microstructure and Cu-precipitation during warm rolling. T h e hardness values could be further improved significantly following a two stage heat treatment process without significantly changing the texture of both the steels. ACKNOWLEDGEMENT T h e authors are greatly acknowledged the help provided by Prof. I. Samajdar of IIT, M u m b a i , India for carrying out the texture measurements in their National Facility of Texture and O I M .
Materials Processing and Texture
171
M i c r o s t r u c t u r e a n d T e x t u r e D e v e l o p m e n t in W a r m R o l l e d H i g h S t r e n g t h IF S t e e l s
REFERENCES ' G . H . Akbari, C M . Sellars, J.J. W h i t e m a n , Acta Metall. Mater. 45 (1997), 5047. M . R . Bamett, J.J. Jonas, ISTJ Int., 39 (1999), 9 , 856. A . Haldar, X. H u a n g . T. Leffers, N. Hensen, R.K. Ray. Acta Mater. 52 (2004), 5 4 0 5 . I . B . Timokhina, A.I. Nosenkov, A . O . H u m p h r e y s , J.J. Jonas, E.V. Pereloma, ISU Int., 44 (2004), 4,717. M . R . Bamett, J.J. Jonas, ISU Int., 37 (1997), 7, 6 9 7 . M . R . Barnett, J.J. Jonas, ISU Int., 37 (1997), 7,706. A . Haldar, R.K. Ray, Mater Sei and Engg. A 391 (2005), 402. 2
3
4
5
6
7
8
M . Sanchez-Araiza, S. Godet, J.J. Jonas, I C O T O M 14. July 11-15, 2 0 0 5 , Leuven, Belgium. H . J . Bunge, Texture Analysis in Material Science, Butterworths, 1982. B . L . Li, A. Godfrey, Q . C M e n g , Q. Liu, N. Hensen, Acta Mater, 52 (2004), 1069. " B . Bay, N. Hansen, D. A. Huges, D. Kuhlman-wilsdorf, Acta Metall. Mater. 40 (1992) 2 0 5 . N . Hansen, D.J. Jensen, Trans R S o c , Lond. A 357 (1999), 1447. H . Regie, A. Miroux, Y. Raulet. L. Chapuis, Proc. of 21st Riso Int. symp. of Mater. Sei. (2000), 525. Q . Liu, B.L. Li, W . Liu, X. Huang, Proc. of 21st Riso Int. s y m p . of Mater. Sei., (2000), 4 2 3 . M . Morita, K. Sato, Y. Hosoya, ISU Int., 34 (1994), 1 92. R . Rana, W. Bleck, S.B. Singh, O.N. Mohanty, Mater. Tech. Carbon Steels, Steel Research Int., 78 (2007), 8, 612. R . Rana, W . Bleck, S.B. Singh, O.N. Mohanty, Mater. Tech. Carbon Steels, Steel Research Int., 78(2007), 8,631. R . Rana, W . Bleck. S.B. Singh, O.N. Mohanty, Mater. Tech. Carbon Steels, Steel Research Int., 78 (2007), 8, 622. 9
10
12
I 3
I 4
1 5
16
I7
18
172
Materials Processing a n d Texture
E F F E C T O F TIN O N T H E T E X T U R E O F N O N O R I E N T E D E L E C T R I C A L S T E E L S < n
C h u n - K a n H o u , T o n g - Han Lee
( l )
a n d Chun-Chih L i a o
(2)
" ' D e p a r t m e n t of Mechanical Engineering, (2)
G r a d u a t e School of Engineering Science and Technology - Doctoral Program, National Yunlin University of Science and Technology, 123 University Road, Section 3 , Douliou, Yunlin 6 4 0 0 2 , Taiwan, R.O.C. ABSTRACT The effect of tin ( range from 0.029- 0.21 w t % ) on the inclusion size in hot bands, grain size and texture of four non-oriented electrical steels after final annealing has been investigated. The steel ingots were reheated to 1200°C and finished rolled at 950 °C. Then hot bands were annealed at 750 °C for t w o hours and cold rolled to 0.5mm. After annealing the cold rolled sheets at 750-950 °C for one minute, both X-ray diffractometer with texture attachment and E B S D analyses were conducted to m e a s u r e specimen texture. It w a s found that 5-10 μηι coarse M n S + A1N inclusions existed in the hot b a n d s containing 0.029 w t % tin. However, 1-2 μηι fine M n S + A I N inclusions were found in the hot bands containing 0.21 w t % tin. After final annealing, grain size decreased with increasing tin content and began to level off when tin reached 0.10wt%. At the same cold rolled annealing temperature, the intensity of {011}<100> texture increased with increasing tin content. In the s a m e steel, the intensity of {011(<100> texture increased with increasing annealing texture and reached a m a x i m u m value and then declined sharply . T h e temperature to obtain m a x i m u m intensity of {011}<100> texture strongly correlated with tin content. E B S D showed that the n u m b e r and volume fraction of {011 }<100> grain per m m increased with increasing tin content. It is concluded that adding tin in the range of 0.100 - 0.210wt% to non-oriented electrical steels can enhance the volume fraction of Goss grain. T h e m e c h a n i s m s are that tin not only enhances the nucleation of Goss grain, but also promotes its growth. Conversely, adding 0.055wt% tin to non-oriented electrical steels can lead to the best cube texture. This tin content maximizes the nucleation density of cube grains during annealing at 850°C. 2
INTRODUCTION Non-oriented electrical steels are important functional materials that are used as lamination cores in various motors and generators. D u e to a general shortage of energy, it is important to improve the efficiency of many electrical machines. Therefore, many researchers are devoted to the study and development of high magnetic induction and low core loss electrical steels. M a n y factors, such as chemical composition, impurities, inclusions, grain size, and texture influence the magnetic properties of non-oriented electrical steels [1-3]. Electrical steel is an anisotropic material in magnetic properties. It has easy magnetization in the <001> crystallographic direction, but it has hard magnetization in the <111> direction. Therefore, {001 }<100> cube texture and {011} < 100> Goss texture are of benefit to the magnetic properties of electrical steels. On the other hand, {111 } fiber textures are deleterious to the magnetic properties of electrical steels. Shimanaka et al. [4-5] reported that adding antimony into electrical steels can enhance the development of {001 } texture. G o d e c et al. [6-7] reported that tin can promote the development of {001 }<100> texture in non-oriented electrical steels. In addition, I w a y a m a et al. [8] studied the effect of copper and tin o n the magnetic properties of
173
Effect of Tin o n t h e T e x t u r e of N o n o r i e n t e d Electrical S t e e l s
grain oriented electrical steels. However, the real m e c h a n i s m s h o w tin affect texture development in electrical steel remain unclear. Recently, we have studied the effects of alloy elements and process variables on the magnetic properties of non-oriented electrical steels [9-11]. This has stimulated us to investigate the effects of tin on the texture development of non-oriented electrical steels. Both orientation distribution function and electron backscatter diffraction ( E B S D ) were conducted to identify texture c o m p o n e n t s in steel sheets. Quantitative results and volume fractions of specific textures were measured to clarify the nucleation and growth mechanisms. EXPERIMENTAL PROCEDURE Four non-oriented electrical steels with various tin contents were prepared in a 250 kg vacuum induction furnace and cast into 4 0 0 m m χ 2 8 0 m m χ 100 m m ingots. Their tin contents were 0.029, 0.055, 0.100, 0.210 w t % . Table I shows their chemical compositions. The steel ingots were reheated to 1200°C for two hours, and hot-rolled in two steps by a two-high pilot hot rolling mill with intermediate annealing at 1200°C for l h o u r . The hot-rolled finishing temperature was 950°C. The thickness of steel ingots after rough rolling was 15mm, and after finishing rolling w a s 2.8 m m . After air cooling to room temperature, hot-rolled bands were annealed at 750°C for 2 hours and furnace cooled. Then, the annealed hot-rolled bands were uniformly milled to 2.3 m m . A four high pilot cold rolling mill was used to cold roll the hotrolled bands to a final thickness of 0.5mm with a total reduction of 7 8 % . Then, the cold-rolled steels sheets were cut into 6 0 m m χ 3 0 m m c o u p o n s both in longitudinal and transverse directions with respect to the rolling direction. Finally, the coupons were annealed in a vacuum infrared furnace set at 750, 800, 850, 900 and 950°C for one minutes with a heating rate 5 °C per second to simulate the continuous annealing. After cold-rolled final annealing, microstructures of both annealed hot-rolled bands and final annealing cold-rolled steel sheets were observed by optical microscope and image analyzer. Grain size of the cold-rolled steel sheets after final annealing was measured by using the intercept method based on A S T M standard E-l 12.
Table I. Chemical compositions of tin added non-oriented electrical steels, (wt % ) Ρ AI Heat N o . Si Mn C Ν ο Sn S A
1.20
0.30
0.013
0.34
0.002
0.005
0.001
0.0021
Β
1.22
0.31
0.014
0.34
0.002
0.005
0.001
0.0014
0.055
C
1.20
0.32
0.013
0.34
0.002
0.004
0.001
0.0028
0.100
D
1.22
0.32
0.013
0.33
0.002
0.004
0.001
0.0024
0.210
0.029
After final annealing, cold-rolled specimens for texture examination were ground to quarter thickness from surface and etched by a hydrogen peroxide solution mixed with several drops of hydrofluoric acid. Each specimen w a s analysed to measure (110), (200), and (112) pole figures by using a S i e m e n s D-5000 X-ray diffractometer with a texture attachment. The diffractometer with M o radiation w a s operated at 40kV and 25 m A . Then three pole figure data were used to calculate the orientation distribution by O D F software with L = 2 0 . Individual grain orientations of annealed hot-rolled bands and final annealed cold-rolled steel sheets were measured using a Hitachi S-30 scanning electron microscope with a H K L E B S D attachment. Then, the n u m b e r and volume fraction of Goss and cube grain were measured.
174
Materials Processing and Texture
Effect of Tin o n t h e T e x t u r e of N o n o r i e n t e d Electrical S t e e l s
RESULTS AND DISCUSSION Microstructures The effect of tin content on inclusion size in the annealing hot-rolled bands is shown in Figures 1-2. Inclusion size decreases with decreasing tin content in the hot-rolled bands of nonoriented electrical steels. T h e size of inclusions in the steel containing 0.029 w t % tin range from 5-10 μηι as s h o w n in Figure 1. T h e types of inclusions in the hot-rolled bands were identified as complex c o m p o u n d s of A l j O j , A1N and M n S . The size of inclusions in the steel containing 0.210 w t % tin range from 1-3 μιτι as s h o w n in Figure 2. The types of inclusions in the hot-rolled bands were identified as simple A1N or c o m p l e x c o m p o u n d s of A I 2 O 3 , A1N and M n S . After final coldrolling sheet annealing, the effect of tin content and annealing temperature on grain size is shown in Figure 3. At the same annealing temperature, grain size decreases with increasing tin content in steel sheets. The grain size levels off for tin content more than 0.1 wt%. In the same steel sheets, grain size increases with increasing annealing temperature. However, during annealing at 750"C, steel sheets with more than 0.055 w t % do not complete recrystallization as shown in later E B S D pictures.
Figure 1. (a) Inclusions in the hot band of steel containing 0.029wt% tin, (b) E D S analysis of inclusions. F.T.=950°C, Ta=750°C.
Figure 2. (a) Inclusions in the hot band of steel containing 0.210wt% tin, (b) E D S analysis of inclusions. F.T.=950"C. Ta=750"C. T a k a s h i m a et al. [3] reported that tin can reduce the inclusion size in grain oriented electrical steels. In this study, the same results are found in the hot rolled bands of non-oriented electrical steels. Tin segregates readily to surfaces, grain boundaries, and defects during hot
Materials Processing and Texture
175
Effect of Tin o n t h e T e x t u r e of N o n o r i e n t e d Electrical S t e e l s
rolling and hot band annealing. It is postulated that tin slows d o w n the nucleation and growth of inclusions near the structural defects. Therefore, the inclusion size decreases with increasing tin content in steels. Since the sulfur and nitrogen content in four electrical steels are almost the same, die total n u m b e r of inclusions in steels decrease with increasing inclusion size. Furthermore, the cold rolling reduction w a s the s a m e for all steels. According to Zener theory, the final grain size is proportional to inclusion size and inversely proportional to v o l u m e fraction of inclusion. So, for the s a m e annealing temperature, grain size decreases with increasing tin content in steel sheets.
Figure 3 . Effects of tin content and final annealing temperature on the grain size. Texture development after final annealing The effect of tin content on the orientation distribution functions of steel sheets after final annealing at 800"C is shown in Figure 4. The predominant orientations in the final annealing sheets are ( O U } < 1 0 0 > Goss texture and [ l l l j < u v w > γ fiber texture. The intensity of G o s s texture increases with increasing tin content in steel sheets. H o w e v e r , the intensity variation of γ fiber texture with tin content is not so regular. Furthermore, {001}<100> cube texture also appears in all steel sheets. T h e intensities of cube texture are nearly the s a m e . At other annealing temperatures, the major texture c o m p o n e n t s are the same. Goss texture also increases with increasing tin content in steel sheets. After annealing at 850"C, the effect of tin content and annealing temperature on the skeleton variations of {001 j < 1 0 0 > , ( O i l J<100> and (111 } texture are shown in Figure 5. G o s s texture h a s relatively stronger intensity than cube and γ fiber textures. Its intensity increases with increasing tin content in steel sheets. H o w e v e r , the intensity différence in cube and γ liber textures is relatively small. The effect of tin content and final annealing temperature on the intensity of {011}<100> is s h o w n in Figure 6. In the steel with the same tin content, there exists a m a x i m u m { 0 1 1 j < 1 0 0 > intensity. T h e o p t i m u m annealing temperature to obtain m a x i m u m {011 }<100> intensity depends upon the tin content in steel sheets. T h e o p t i m u m annealing temperature increases with increasing tin content and reaches a m a x i m u m at 900"C, then, it drops.
176
Materials Processing a n d Texture
Effect of Tin o n t h e T e x t u r e of N o n o r i e n t e d Electrical S t e e l s
Figure 4. Effect of tin content o n the orientation distribution function of steels after final annealing at 800 °C.
Figure 5. Effect of tin content on the variation of texture intensity after annealing at 850 °C. , (a) { 0 0 1 } < 1 0 0 > a n d {011}<100>. (b) [111 J < u v w > .
Materials Processing and Texture
177
Effect of Tin o n t h e T e x t u r e of N o n o r i e n t e d Electrical S t e e l s
Figure 6. Effect of tin content and final annealing temperature o n the intensity of |0!1|<100>. T h e orientation images of the steels with various tin contents after final annealing at 750 °C and 850 °C are s h o w n in figures 7-8, respectively. In Figure 7, deformed band structures still exists in the steels containing m o r e than 0.055 w 1 % tin. T h e m a i n orientation of the deformed band structures is {100}<0l l > . Furthermore, a lot of { I 0 0 } < 0 0 l > texture occurs in the central portion of the steel sheets with 0.055 w t % tin. However, many (011 }<100> grains locate at the quarter thickness layer from the surface in the steel with 0.210 w t % tin. In Figure 8. all the steels complete recrystallization and grains b e c o m e coarse. T h e effects of tin content and annealing temperature on the {011 }<100> grain n u m b e r are s h o w n in Figure 9. It is found that the n u m b e r of Goss grain increases with increasing tin content in the steel sheets. T h e n u m b e r of Goss grain decreases with increasing annealing temperature. The effects of tin content and annealing temperature on the volume fraction of [011 }<100> grain are shown in Figure 10. The v o l u m e fraction of Goss grain increases with increasing tin content in the steel sheets. In the steel with more than 0.100 w t % tin, v o l u m e fraction increases with increasing annealing temperature. The effects of tin content and annealing temperature on the {001 }<100> grain n u m b e r is s h o w n in Figure 11. During annealing at 850°C. the n u m b e r of cube grains increases with increasing tin and reaches a m a x i m u m value in the steel with 0.055 w t % tin. Thereafter, it drops and the curve becomes flat. H o w e v e r , the n u m b e r density o f cube grains remains at almost at the same level during annealing at 750°C. The effects of tin content and annealing temperature o n the v o l u m e fraction of {001 ) < 1 0 0 > grain are shown in Figure 12. During annealing at 850"C, the steel with 0.055 w t % tin has m a x i m u m volume fraction o f cube grain. A s the results mentioned above, according to Figures 4-6 and 9-10. it was found that adding tin in the range of 0.100 - 0.210 w t % in non-oriented electrical steels can enhance the volume fraction of Goss grain. T h e m e c h a n i s m s are that tin is not only to e n h a n c e s the nucleation of G o s s grain, but also promotes the growth of Goss grain. O n the other hand, the best cube texture occurred in those non-oriented electrical steels with 0.055 w t % tin. This tin content maximizes the nucleation density of cube grains during annealing at 850°C.
178
Materials Processing and Texture
Effect of Tin o n t h e T e x t u r e of N o n o r i e n t e d Electrical S t e e l s
Figure 7. Orientation images of steels with various tin content after final annealing at 750 °C. (a) 0.029wt% Sn, (b) 0 . 0 5 5 w t % Sn, (c) 0.100 w t % Sn. (d) 0.210 w t % Sn.
Materials Processing and Texture
179
Effect of Tin o n t h e T e x t u r e of N o n o r i e n t e d Electrical S t e e l s
Figure 8. Orientation images of steels with various tin content after final annealing at 850 °C. (a) 0.029wt% Sn, (b) 0 . 0 5 5 w t % Sn, (c) 0.100 w t % Sn, (d) 0.210 w t % Sn.
180
Materials Processing and Texture
Effect of T i n o n t h e T e x t u r e of N o n o r i e n t e d Electrical S t e e l s
Figure 9. Effect of tin content and annealing temperature on the {011 }<100> grain number.
Figure 10. Effect of tin content and annealing temperature on the volume fraction of {011}<100> orientation grains.
Figure 11. Effect of tin content and annealing temperature on the {001 }<100> grain number.
Figure 12. Effect of tin content and annealing temperature on the volume fraction of {001 }<100> orientation grains.
CONCLUSION The effects of tin on the texture development of non-oriented electrical steels has been investigated. These are summarized as follows: 1. In the annealing hot-rolled bands, inclusion size decreases with decreasing tin content in the hot-rolled bands of non-oriented electrical steels. With annealing temperature kept constant, grain size decreases with increasing tin content in the steel sheets. It does not change, however.
Materials Processing and Texture
181
Effect of Tin o n t h e T e x t u r e of N o n o r i e n t e d Electrical S t e e l s
for tin contents more than 1 wt%. In the same steel sheets, grain size increases with increasing annealing temperature. It is postulated that tin can reduce and retard the nucleation and growth of inclusions near the structural defects. 2. The predominant orientations in the final annealing sheets are {011 }<100> Goss texture and {111 } γ fiber and cube textures. The intensity of Goss texture increases with increasing tin content in steel sheets. It is found that adding tin in the range of 0.100 - 0.210 w t % in nonoriented electrical steels can enhance the volume fraction of G o s s grain. T h e m e c h a n i s m s are that tin not only to enhances the nucleation of Goss grain, but also promotes the growth of Goss grain. Conversely, the best cube texture occurred in the non-oriented electrical steels with 0.055wt% tin. This tin content m a x i m i z e s the nucleation density of cube grains during annealing at 850°C. ACKNOWLEDGMENT The authors would like to appreciate the financial assistance of National Science Council of Taiwan through the grant n u m b e r N S C 9 4 - 2 2 1 6 - E - 2 2 4 - 0 0 7 . REFERENCES Έ . J. Hillinski, "Recent developments in semiprocessed cold rolled magnetic lamination steel", J. Magn. M a g n . Mater., 304, 172-177 (2006). 2
E . T. Stephenson and M.R. A m a n n , "Effects of Composition, Grain Size and Texture on the A C Core Loss of magnetic Lamination Steel", in Energy Efficient Electrical Steels, A.R. Marder and ET. Stephenson, Ed., T M S - A I M E , 43-60 (1981). 3
M . Takashima, N. M o r i t o ; " Nonoriented Electrical Steel Sheet with L o w Iron Loss for High-Efficiency Motor C o r e s " , I E E E Transactions on Magnetics, 3 5 , 557-561 (1999). 4
H . Shimanaka, Y. Ito, T. Irie, K. Matsumura, H. N a k a m u r a , and Y. Shono, "Non-oriented Si-Steels Useful For Energy Efficient Electrical A p p a r a t u s " , in Energy Efficient Electrical Steels, A. R. Marder and Ε. T. Stephenson, Eds. T M S - A I M E , 193-204 (1981). 5
H . Shimanaka, T. Irie, K. M a t s u m u r a and H. N a k a m u r a , " A N e w Non-oriented Si-Steel with Texture of { 1 0 0 } < 0 v w > " , J. M a g n . M a g n . Mater., 19, 63-64 (1980). °M. Godec, M. Jenko, H. J. G r a b k e and R. Mast, "Sn Segregation and Its Influence on Electrical Steel Texture Development", ISU International, 39, 742-746 (1999). 7
M . Godec, M . Jenko, R. Mast, and H. J. G r a b k e " Texture measurements on electrical steels alloyed with tin", V a c u u m , 6 1 , 151-155 (2001). 8
K . Iwayama, K. Kuroki, Y. Yoshitomi, and K. H o m m a , T. W a d a , " Roles of tin and copper in the 0.23-mm-thick high permeability grain-oriented silicon steel", J. Appl. Phys. 5 5 , 21362138 (1984). 9
C h u n - K a n Hou. Chen-Ti Hu and Sanboh Lee, " T h e Effect of A l u m i n u m on the Magnetic Properties of Lamination Steels", I E E E Trans. Magn., 2 7 , 4 3 0 5 - 4 3 0 9 (1991). 10
C h u n - K a n Hou. "Effect of Hot Band Annealing Temperature on the Magnetic Properties of L o w - C a r b o n Electrical Steels", ISU International, 36, 563-571 (1996). ' ' C h u n - K a n H o u , "The Effects of Grain Size on the Magnetic Properties of Fully Processed, Continuous-Annealing Low-Carbon Electrical Steels", IEEE Trans. Magn., 32, 471-477 (1996).
182
Materials Processing a n d Texture
E F F E C T O F C O L D R O L L E D R E D U C T I O N O N T H E T E X T U R E A N D PLASTIC STRAIN R A T I O O F 2205 D U P L E X S T A I N L E S S S T E E L Chun-Kan Hou and C h a n g - C h i a Liu Department of Mechanical Engineering, National Yunlin University of Science and Technology, 123 University Road, Section 3 , Douliu, Yunlin 640, Taiwan, Republic of China.
ABSTRACT The effect of cold rolled reduction, range from 52.4 % to 67 %, and annealing temperature from 1025 - 1100 °C o n the annealing texture and plastic strain ratio of 2205 duplex stainless steel have been studied. The predominant δ phase texture components in cold rolled state were {001 } < 1 1 0 > and {111}<1 12>. Both texture intensities increased with increasing cold rolled reduction and reached a m a x i m u m value at cold rolled reduction 6 1 . 9 % and then became level off. The predominant γ phase texture c o m p o n e n t s in cold rolled state were {011}<311> and { 1 1 2 } < 1 1 1 > . T h e intensity of {011 }<311> increased with increasing cold rolled reduction. After cold rolled sheet annealing, the predominant δ phase texture components were { 0 0 1 } < 1 Î 0 > , {011 } < 0 Î 1>, and { 1 1 1 } < Î Î 2 > . The intensity of { 0 0 1 } < 1 Ï 0 > increased with increasing cold rolled reduction. T h e effect of temperature on texture was not obvious. The predominant γ phase texture c o m p o n e n t s after annealing were {011}<311> and {113}<3 3 2 > . Both intensities of {011}<3 1 1> and {113}< 3 3 2> increased with increasing cold rolled reduction. Furthermore, the intensity of {113} < 3 3 2> increased with increasing annealing temperature in the specimens with cold rolled reduction larger than 6 1 . 9 % . After cold rolled sheet annealing, r value increased with increasing cold rolled reduction. However, r value decreased with annealing temperature and reached a m i n i m u m at 1075 °C. Besides, plane anisotropy absolute Ar value increases with increasing cold rolled reduction. Moreover, absolute Ar value increases with increasing annealing temperature in the specimens with cold rolled reduction less than 5 7 % . Conversely, absolute Ar value decreases with increasing temperature and reached a m i n i m u m at 1075 °C in the specimen with 6 7 % cold rolled reduction. INTRODUCTION Duplex stainless steels have the combined advantages of excellent resistance to chloride stress corrosion cracking, non-sensitive to intergranular corrosion cracking, high strength, low nickel content and cheaper price. Therefore, duplex stainless steels have great potential to substitute austenitic stainless steels in chemical and petrochemical industries, offshore facilities, p o w e r generation cooling system, marine transportation tankers, food industry containers [1]. Generally, duplex stainless steels contain about 30 - 70 volume percentage of ferrite phase depending on their chemical compositions and subsequent heat treatments. For example, the volume percentage of ferrite is almost equal to that of austenite in 2205 duplex stainless steels. Duplex plates and cold rolled sheets are c o m m o n products. During manufacture of flat plates and sheets, grain structures change and tend to deform along the rolling direction. Therefore, the microstructures of duplex stainless steels look like fiber composite material. Texture and anisotropy of properties are important issues for flat cold rolled duplex stainless steels. They affect the subsequent metalworking of duplex stainless steels, such as sheet drawing and forming. Several researches studied the texture developments of austenitic, ferritic, and duplex
183
Effect of C o l d R o l l e d R e d u c t i o n o n t h e T e x t u r e a n d Plastic S t r a i n of 2 2 0 5 D u p l e x S t e e l
stainless steels after hot rolling, cold rolling and annealing [2-10]. Akdut et al. found that the {011} <211 > and {011} < 100> were the m a i n cold rolled texture in the austenite and {112} < 110> and {111 } w e r e the predominant cold rolled textures in the ferrite of high nitrogen duplex stainless steels [2]. Hutchinson et al. found that {110}<112> was the main annealing texture in austenite and {100}<011> was the main annealing texture in ferrite of 2205 duplex stainless steel [6]. Song and Bate found that the m e c h a n i s m s of plastic anisotropy deformation of duplex stainless steel w a s intragranular slip and texture strengthening [7]. However, there are few researches dealing with texture and plastic strain ratio of cold rolled duplex stainless steels [6-7]. The purposes of this study are to investigate the effects of cold rolled reduction and annealing temperature on the development of cold rolled texture, annealing texture, and plastic strain ratio of 2205 duplex stainless steel. The relationships between texture and plastic strain ratio are also explained. EXPERIMENTAL PROCEDURE The starting material is a commercial 2205 duplex stainless steel billet in size of 145 m m (ND) χ 180 m m (TD) χ 230 m m (RD). Its chemical composition is shown in Table 1. The billet is manufactured by electrical arc furnace melting, ladle refining, ingot casting, and v a c u u m arc refining and forged. Hot rolling is conducted by a two high pilot rolling mill in three steps. At first and second steps, the starting billet is reheating t o l 2 0 0 ° C and soaking 1 h for hot rolling and the width is kept as 230 m m . In the first step, the thickness is reduced from 145 m m to 20 m m . After second step, the thickness is reduced to 5 m m . In the third step, the hot rolled plates are reheating to 1250°C and soaking l h . The hot plates are hot rolled to to 2.75, 2.45, 2.33, and 2.11 m m in thickness respectively. Then the hot rolled plates are cut into 500 m m length plates along the rolling direction and annealed at 1 0 3 0 ° C for 0.5 h and air cooling. After that, the annealed hot rolled bands are shot peening by sand p o w d e r s to remove the surface scale. A pilot four high cold mill is conducted to cold rolled the hot plates to 1 m m thickness. The cold rolled reduction ratios are 52.4, 57.0, 61.9, and 6 7 % respectively. Then, the cold rolled sheets are cut into 1 m m χ 30 m m χ 180 m m c o u p o n s along the direction from 0, 4 5 , and 90° to the rolling direction. Table I. Chemical composition of 2205 duplex stainless steel. (wt%) Μη Steel Fe C Si Ρ S Cr Ni 2205 Bal. 0.02 0.55 1.40 0.028 0.001 22.60 4.9
Mo 3.2
Ν 0.20
The cold rolled coupons are annealed by heating up from room temperature to 1025, 1050, 1075, and 1 1 0 0 ° C , respectively and soaking 10 minutes and with a heating rate of 5°C/min in an sealed Lindberg S T F 5 4 4 5 9 C electric tube furnace full with argon atmosphere. At the end of soaking time, annealed steel sheets are taken out and cool down in air. After cold rolled sheets annealing, the plastic strain ratio, r value, along the direction from 0, 4 5 , 90° in the rolling direction are measured according to A S T M E517-96a. T h e n average f value and plane anisotropic Ar value are calculated by following equation.
(1) (2)
184
Materials Processing and Texture
Effect of C o l d Rolled R e d u c t i o n o n t h e T e x t u r e a n d Plastic Strain of 2 2 0 5 D u p l e x S t e e l
S p e c i m e n s at various stages are polished and etched by Vilella's reagent and observed by an O l y m p u s B H 2 - U M A optical microscope. Leco2001 image analysis system is conducted to measure the v o l u m e percentages of ferrite and austenite phases. Texture of both as cold rolled and annealed cold rolled specimens are measured by a Siemens D5000 X-ray diffractometer with a texture attachment. M o target is used and the operating condition is set at 4 0 K V and 35mA. To measure texture, specimens are ground and polished out quarter thickness and etched in a hydrogen peroxide solution added with several drops of fluoride acid. For the texture of ferrite phase, three pole figures including (200), (211). and (222) pole figures are measured. On the other hand, the texture of austenite phase, three pole figures including (200). (220). and (311) pole figures are measured. Three pole figure data are used to calculate the orientation distribution function with L = 2 0 . Counter levels set at 1 . 2 , 4 . 6, 8, 12, and 14 respectively. RESULTS Microstructures After cold rolled at various reductions, the as cold rolled microstructures are shown in figure 1. In the specimens, elongated γ austenite phases lie along the rolling direction in the δ ferrite phases. In the thickness direction, the width of austenitic phase decreases with increasing cold rolled reduction. Furthermore, the distance between austenitic fibers also decreases with increasing cold rolled reduction. After annealing at 1025 - 1100°C for 10 minutes, the microstructures of the specimens with 6 7 % cold rolled reduction are shown in figure 2. Besides the irregular elongated austenite phases, some isolated austenitic islands appear. The width and inter fiber distance also increase with the increasing annealing temperature.
Figure 1. Microstructures of 2205 duplex stainless steel with different cold rolled reduction (C.R.), (a) C.R.=67%. (b) C R = 6 1 . 9 % , (c) C . R . = 5 7 % , (d) C.R.=52.4%.
Figure 2. After 6 7 % cold rolled reduction, microstructure of 2205 duplex stainless steel annealed at, (a) 1025"C. (b) 1050"C. (c) 1 0 7 5 T , (d) 1100"C, for 10 min.
Cold rolled texture T h e effect of cold rolled reduction on the orientation distribution functions of δ ferrite are shown in figure 3. T h e predominant textures are {001} with peak at {001}<110>, [ l l l } < u v w > with peak at {111}<112>. Both intensities of {001}<110> and [111}<112> textures increase with increasing cold rolled reduction and reach a m a x i m u m value at 6 1 . 9 % cold
Materials Processing and Texture
185
Effect of C o l d Rolled R e d u c t i o n o n t h e T e x t u r e a n d Plastic Strain of 2 2 0 5 D u p l e x S t e e l
reduction. Then, they b e c o m e level off. Furthermore, a w e a k } 110} texture with a peak texture at {110}<001> are also developed. The intensity of {110}<001> increases with cold rolled reduction and reaches a m a x i m u m value at 5 7 % cold rolled reduction and then fall down. The effect of cold rolled reduction on the orientation distribution functions of γ austenite are shown in figure 4. T h e predominant texture components are {011}<100>, {011}<311>, [ 0 1 1 j < 2 1 1 > and {112}<111>. The intensity of {011J<311> texture increases with increasing cold rolled reduction. While cold rolled reduction is m o r e than 6 1 . 9 % , the intensity of {011 ) < 3 1 1 > rapidly increases. However, the intensity of {112}<111> texture slightly decreases with increasing cold rolled reduction.
Figure 3 . Orientation distribution function of ferrite in duplex stainless steel with, (a) C R . = 6 7 % , (b) C . R . = 6 1 . 9 % , (c) C . R . = 5 7 % , (d) C.R.=52.4%. (e) Ideal orientation in Euler space with2 = 0 and<J> = 4 5 . 2
Figure 4. Orientation distribution function of austenite in duplex stainless steel with, (a) C . R . = 6 7 % . (b) C R " 6 1 . 9 % . (c) C . R . = 5 7 % . (d) C R " 5 2 . 4 % . Annealing texture After annealing at various temperatures, the orientation distribution functions of ferrite phase are shown in figure 5. The predominant texture c o m p o n e n t s are {001} fiber texture with peak at {001 }<110> and {111 ) < u v w > fiber texture with peak at {111 }<112>. In addition, [011 }<011> texture also appears. The effects of cold rolled reduction and annealing temperature on the intensities of {001}<110> are shown in figure 6. At the same annealing temperature, besides the fluctuation of 6 1 . 9 % cold rolled reduction, the intensity of { 0 0 1 j < 1 1 0 > texture in ferrite increases with increasing cold rolled reduction. The effect of cold rolled reduction and
186
Materials Processing and Texture
Effect of C o l d Rolled R e d u c t i o n o n t h e T e x t u r e a n d Plastic Strain of 2 2 0 5 D u p l e x S t e e l
annealing temperature on the intensity of {111}<112> are s h o w n in figure 7. Despite the specimens with 6 1 . 9 % cold rolled reduction, the intensity of {111}<-1-12> increases with increasing cold rolled reduction after annealing at the same temperature. While cold rolled reduction is larger than 5 7 % . the peak intensities of {111}<-1-12> texture occurs at 1075 1100"C.
Figure 5. After 6 7 % cold rolled reduction and annealing at different temperature for 10 min., orientation distribution function of ferrite in duplex stainless steel annealed at, (a) 1 0 2 5 T . (b) 1050°C, (c) 1075°C, (d) 1100°C. After annealing at various temperatures, the orientation distribution functions of austenite phase are s h o w n in figure 8. The m a i n texture c o m p o n e n t s are [011 }<-311> and {113}<-3-32>. The effects of cold rolled reduction and annealing temperature o n the intensity of {011}<-311> are shown in figure 9. At the s a m e annealing temperature, specimens with cold rolled reduction larger than 6 1 . 9 % h a v e m o r e stronger intensity of {011}<-311> than that with smaller cold rolled reduction. T h e effects of cold rolled reduction and annealing temperature on the intensity of {113}<-3-32> are shown in figure 10. At the s a m e annealing temperature, intensity of {113}<-3-32> increases with increasing cold rolled reduction. If cold rolled reduction is larger than 6 1 . 9 % , intensity of {113}<-3-32> increases with increasing annealing temperature.
Figure 6. Effects of cold rolled reduction and annealing temperature on intensity of [001 }<110> in ferrite phase.
Figure 7. Effects of cold rolled reduction and annealing temperature on intensity of {111 [ < 1 1 2 > in ferrite phase.
Materials Processing and Texture
187
Effect of C o l d Rolled R e d u c t i o n o n t h e T e x t u r e a n d Plastic Strain of 2 2 0 5 D u p l e x S t e e l
Figure 8. After 6 7 % cold rolled reduction and annealing at different temperature for 10 min., orientation distribution function of austenite in duplex stainless steel annealed at, (a) 1025"C, (b) 1050"C, (c) 1075"C. (d) 1100"C. Plastic strain ratio The effect of cold rolled reduction and annealing temperature on the plastic strain ratio r value is shown in figure 1 1 . At the s a m e annealing temperature, r value increases with increasing cold rolled reduction. However, at the same cold rolled reduction, f value changes with annealing temperature. Besides the specimens with 6 1 . 9 % cold rolled reduction, there is a m i n i m u m r value occurs at 1075"C. The effect of cold rolled reduction and annealing temperature on the plane anisotropy of duplex stainless steel, Ar value, are shown in figure 12. The Ar value is between -0.65 to -1.25. At the same annealing temperature, absolute Ar value increases with increasing cold rolled reduction. Despite the specimens with 6 7 % cold rolled reduction, absolute Ar value increases with increasing annealing temperature. On the contrary, absolute Ar value of specimens with 6 7 % cold reduction decreases with increasing annealing temperature and reaches a m i n i m u m value at 1075"C.
Figure 9. Effects of cold rolled reduction and annealing temperature on intensity of {011 }<311> in austenite phase.
188
Materials Processing and Texture
Figure 10. Effects of cold rolled reduction and annealing temperature on intensity of {113 j < 3 3 2 > in austenite phase.
Effect of C o l d Rolled R e d u c t i o n o n t h e T e x t u r e a n d Plastic Strain of 2 2 0 5 D u p l e x Steel
Annealing temperature ( T )
Figure 11. Effects of cold rolled reduction and annealing temperature on f value of duplex stainless steel.
Annealing temperature (°C)
Figure 12. Effects of cold rolled reduction and annealing temperature on Δ Γ value of duplex stainless steel.
DISCUSSION Texture development M a n y factors such as dislocation slip, twinning, grain boundary sliding, degree of cold reduction, and microstructures can affect the cold rolling reduction. In this study, annealing hot b a n d s of 2205 duplex stainless steels contain ferrite and austenitic phases. The volume fractions of ferrite and austenitic phases are 60 and 4 0 % respectively. And the Vickers microhardness of austenite phase is a little higher than ferrite phase. Furthermore, the range of cold rolling reduction is from 52.4 to 67%.Therefore. dislocation slip is the predominant deformation m e c h a n i s m for both phases. The possible active slip systems in ferrite phase are {110}<111>. {112}<111> and {123}<111>. Raabe [12] simulated rolling textures of b. c. c. metal by grain interaction m o d e l s . H e found that {001}<110>. {112}<110> and {111}<110> α fiber texture and {111 }<110} and {111 }<112} γ fiber texture developed in the pencil glide condition. In this study, the m a i n textures in ferrite phase are {001} α fiber with peak at (001 )[ 110] and { l l l } < u v w > γ fiber with peak at (111)[112]. It implies that {110}, {112} and {123} planes are active during rolling. Moreover, the increase in {111 }<112> intensity can be explained by the relaxed constrains model with additional shear in rolling direction [2]. Austenitic stainless steel has low stacking energy and likes to develop brass-type rolling texture. It is a {110}<112> texture and spread to G o s s type texturef 110}<001>. In this study, the predominant rolling textures in austenite are { 0 U } < 1 0 0 > . {011}<311>. {011}<211> and {112}<111>. Therefore, it is a brass-type texture. But the m a x i m u m intensity is at {011 }<311>. Furthermore, its intensity rapidly increases after cold reduction reaches 6 1 . 9 % . Dillamore et al. [13] found that {110 j < 112>. {110} <001 > and near {4,4.11} < 11.11,8> are stable orientation in f. c. c. crystals after rolling. In this study, the m a x i m u m cold reduction is 6 7 % , it m a y not attain the reduction need to reach final stable orientation. Moreover, {112}<111> decreases with increasing cold reduction. It is an intermediate texture during cold reduction. Finally, the twinning of {112}<111 > texture rotates to Goss orientation [2]. The recrystallization textures of ferrite phase c o m p o s e of {001 } fiber texture with m a x i m u m at {001}<110>. and {001 }<010> and { l l l } < u v w > with peak at {111}<110} and
Materials Processing and Texture
189
Effect of C o l d Rolled R e d u c t i o n o n t h e T e x t u r e a n d Plastic Strain of 2 2 0 5 D u p l e x S t e e l
{001}<110> to {111}<110>, and {110}<110> to {110}<111> orientation. Lewis et al. [8] pointed that {100} grains parallel to the rolling plane nucleated at the adjacent matrix-deformed grain in the ferritic stainless steels. Furthermore, {110} grains parallel to rolling plane nucleated at deformation bands and {111} grains parallel to rolling plane nucleated at grain boundary. In this study, the recrystallization texture of {100}, {110} and {111} grains parallel to rolling plane all occurred. Moreover, cold reduction enhances the intensity of {001 }< 110> and {111 }<112>. It implied that cold reduction can increase the strain energy at adjacent matrix-deformed grain and grain boundary. Furthermore studies on the microstructures are needed to clarify the cold reduction on the nucleation and growth of ferrite grain during annealing. The main recrystallization textures of austenite phase are {011 }<311> and {113}<3-32>. Their intensities are relative weak. T h e {011}<311> texture is very close to {110}<112> brass type texture. Donadille et al. [9] found that {110}<225> texture with a spread toward the {110}<001> were the predominant texture in the annealed austenitic stainless steel. They attributed this m e c h a n i s m to the oriented nucleation in the shear bands of {110}<001> grains invading the {111 } oriented d o m a i n s without nucleation. The {113}<332> texture has not reported yet in duplex stainless steel. Donadille et al. [9] pointed that {113}<211> recrystaalization texture related to the prominent deformation texture c o m p o n e n t s of brass type by rotation of 2 5 - 4 0 ° about a c o m m o m < 1 1 1 > axis. It is postulated the same mechanism occurred in the development of {113}<332> recrystallization texture. Relationship between plastic strain ratio and texture Plastic strain ratio and plane anisotropy are close related to texture parameters. For ferritic stainless steels, {111} and {332} planes parallel to rolling plane are benefit to f value. However, {100} and {112} planes parallel to rolling plane are deleterious to f value. Furthermore, plane anisotropy Δ Γ value decreases with increasing intensity of {112} plane parallel to rolling plane [11]. For austenitic stainless steel, r value is relative low. Its Δ Γ value is related to the difference between intensity of {100} plane and intensity of {110} and {112} planes. For duplex stainless steel, Moverare et al. [14] found that the load partitioning between the two phases is dependent on the loading direction. They also found that r value for the austenite phase was close to the isotropic value ( r =1). H o w e v e r , the ferrite s h o w s strong anisotropy, with low r values in the rolling direction. But, m a x i m u m r value in the surface layer occurred in the transverse direction. Furthermore, m a x i m u m r value in the center layer occurred in the 45-deg direction. In this study, the volume fraction of ferrite phase is about 6 0 % . As mentioned earlier, the intensity of {111 } fiber texture in ferrite increases with increasing cold reduction. So, the r v a l u e also increases with increasing cold reduction. However, the intensity of {112} plane parallel to rolling direction in the ferritic phase is very low. It is postulated that the texture gradient in ferrite phase affects Δ Γ value. Their quantitative relationships need further analysis. CONCLUSION The effects of cold rolled reduction on the texture and plastic strain ratio have been studied. T h e summaries are as follow. 1. The predominant cold rolled textures in ferrite are {001} fiber with peak at {001}<110> and { l l l } < u v w > fiber with peak at {111}<112>. Both intensities of (001 )[ 110] and (111)[112] textures increase with increasing cold rolled reduction and reach a m a x i m u m value at 6 1 . 9 % cold reduction. Then, they b e c o m e level off. The main textures in
190
Materials Processing a n d Texture
Effect of C o l d R o l l e d R e d u c t i o n o n t h e T e x t u r e a n d Plastic Strain of 2 2 0 5 D u p l e x S t e e l
austenitic phase are {011 } ζ fiber texture with peak at {011 }<311>, and {112}<111>. The intensity of {011}<311> texture increases with increasing cold rolled reduction. H o w e v e r , the intensity of (112)[111] texture slightly decreases with increasing cold rolled reduction. The texture developments are consistent to R C Taylor model. 2. The predominant recrystallization textures in ferrite are {001 } fiber texture with peak at {001}<110> and { l l l } < u v w > fiber texture with peak at {111}<112>. At the same annealing temperature, the intensity of {001}<110> texture in ferrite increases with increasing cold rolled reduction. Despite the specimens with 6 1 . 9 % cold rolled reduction, the intensity of {111 }<112> increases with increasing cold rolled reduction after annealing at the s a m e temperature. The main recrystallization textures in austenite are {011 }<311> and {113}<332>. At the same annealing temperature, specimens with cold rolled reduction larger than 6 1 . 9 % have more stronger intensity of {011}<311> than that with smaller cold rolled reduction. If cold rolled reduction is larger than 6 1 . 9 % , intensity of {113}<332> increases with increasing annealing temperature. 3. At the same annealing temperature, r-value increases with increasing cold rolled reduction. H o w e v e r , at the same cold rolled reduction, m i n i m u m f - v a l u e was occurred dependent on annealing temperature. At the same annealing temperature, absolute Δ Γ value increases with increasing cold rolled reduction. Despite the specimens with 6 7 % cold rolled reduction, absolute ΔΓ value increases with increasing annealing temperature. Acknowledgment W e appreciate the kind assistance of Gloria Material Technology Corporation to provide the materials. REFERENCES R. N . G u n n , " Duplex stainless steels: Microstructure, Properties and Application," Abington Publishing, Cambridge, England, 175-185 (1997) N. Akdut and J. Foct, "Phase Boundary and Deformation in High Nitrogen Duplex Stainless Steels I. Rolling Texture D e v e l o p m e n t , " Scripta Metall, et Mater., 32, 103-108 (1995). Ν . A k d u t and J. Foct, "Phase Boundary and Deformation in High Nitrogen Duplex Stainless Steels II. Analysis of Deformation M e c h a n i s m s by Texture Measurements in X 2 C r N i M o 22 5(1.4462)," Scripta Metall, et Mater., 32, 109-114 (1995). A. ul-Haq, Η. Weiland, and Η. J. B u n g e , " Textures and microstructures in duplex stainless steel," Mater. Sei. and Tech., 10, 289-298 (1994) J.A. Jimenez, M. Carsi, O.A. R u a n o , " Characterization of a δ/γ duplex stainless steel," J. Mater. Sei., 3 5 , 9 0 7 - 9 1 5 ( 2 0 0 0 ) W. B. Hutchinson, K. Ushioda, G. Runnsjo, " Anisotropy of tensile behavior in a duplex stainless steel sheet," Mater. Sei. and Tech., 1, 728-731 (1985) J. L. Song and P. S. Bate, " Plastic Anisotropy in a Superplastic Duplex Stainless Steel," Acta Mater., 4 5 , 2 7 4 7 - 2 7 5 7 (1997) D. B. Lewis, F.B. Pickering, " Development of recrystallization textures in ferritic stainless steels and their relationship to formability," Metal Tech., 10, 264-273 (1983) C. Donadille, R. Valle, P. Dervin and R. Penelle, " D e v e l o p m e n t of Texture and Microstructure During Cold-Rolling and Annealing of F.C.C. Alloys: Example of a Austenitic Stainless Steel," Acta Metall., 37, 1547-1571 (1989) 1
2
3
4
5
6
7
8
9
Materials Processing and Texture
·
191
Effect of C o l d Rolled R e d u c t i o n o n t h e T e x t u r e a n d Plastic Strain of 2 2 0 5 D u p l e x S t e e l
H. Honneff, H. Mecking, " Analysis of the Deformation Texture at Different Rolling Conditions," Proc. I C O T O M 6, T o k y o , in "Textures of M e t a l s , " ed. S N a g a s h i m a , The Iron and Steel Inst, of Japan, p.347-355 (1981) " M. Yantac et al, " Texture Development in Ferritic Stainless Steel Sheet," Texture, 1, 71-86 (1972) D. Raabe, "Simulation of rolling texture of b.c.c. metals containing grain interactions and crystallographic slip on {110}, {112} and {123} p l a n e s , " Mater. Sei. Eng., A 1 9 7 , 31-37 (1995) 1 . D. Dillamore, E. Butler, and D. Green, " Crystal rotations under conditions of imposed strain and the influence of twinning and cross slip," Metal Sei., 2, 161-167 (1968) J. J. Moverare and M. Oden, " Influence of Elastic and Plastic Anisotropy on the Flow Behavior in a Duplex Stainless Steel," Metall. Mater. Trans. A, 33A, 57-71 (2002) 1 2
1 3
14
192
·
Materials Processing and Texture
I M P O R T A N C E O F D I S L O C A T I O N S U B S T R U C T U R E S IN T E X T U R E C O N T R O L O F L O W CARBON STEEL Hirosuke Inagaki ΓΝΑΤΕΧ Research Center, Kugenuma-Fuj igaya 4-1-19, Fujisawa-shi, Japan 251 -0031
ABSTRACT Hitherto, the mechanism of the texture formation has been discussed mainly in terms of Taylor's h o m o g e n e o u s deformation theory. However, in order to achieve more precise and effective texture control, the m e c h a n i s m should be investigated in a much finer scale, i. e., at the level of dislocations. In this paper, the development of dislocation substructures during cold rolling was investigated in detail in high purity iron by using T E M and the contribution of these dislocation substructures to the formation of rolling and recrystallization textures w a s evaluated. It was found that so called microband structures, which were developed during cold rolling, were not formed by the formation of isolated microbands and their clustering. It turned out that they were formed by the disintegration of equi-axed cells and their subsequent rearrangement into elongated cells. The development of cold rolling textures seems to be controlled by the development of these micro band structures, since the thickness of the microband limits the mean free path of moving dislocations inducing slip rotation. On subsequent annealing, these microband structures provided new recrystallized grains. Therefore, control of these microband structures is very important in the texture control of iron and steel. INTRODUCTION Texture control is especially important in iron and steel, since most of their important mechanical properties are strongly influenced by textures introduced during fabrication. Also in recently developed n e w materials, texture control plays a very important role. It has been found that, strength , toughness and de-lamination b e h a v i o r ' in high strength pipe-line steels, and ultrasonic sound velocities in very thick high strength steel plates for ship b u i l d i n g are strongly influenced by textures developed during modern thermo-mechanical treatments. Proper control of these textures is still a t h e m e of current investigations. 1
1
2
3
4
In the past, as far as the practical application is concerned, texture control has been most successful in low carbon steels for car body panels. The most representative example will be found in IF steels, in which very high Lankford's r value above 2.2 can b e readily obtained. However, although many investigations have been made so far on the formation of rolling and recrystallization textures in iron and low carbon steels, its detailed m e c h a n i s m is not clear y e t ' Since recrystallization textures in iron and steel are most significantly affected by cold rolling textures, control of cold rolling textures is very important in their texture control. These cold rolling textures are developed due to slip rotation which is induced by glide of moving dislocations. At high rolling reductions where cold rolling textures are developed remarkably, 5
7
193
I m p o r t a n c e of D i s l o c a t i o n S u b s t r u c t u r e s in T e x t u r e C o n t r o l of L o w C a r b o n S t e e l
these dislocations are present in the form of dislocation substructures. Therefore, control of these dislocation substructures is m o s t essential in the control of cold rolling textures. Therefore, in this paper, the development of dislocation substructures during cold rolling was investigated in detail by using high purity iron. The results obtained were compared with the development of the global cold rolling textures revealed with the X ray O D F (orientation distribution function) analysis, and the contribution of the dislocation substructures to the development of cold rolling textures w a s evaluated. EXPERIMENTAL The starting materials were hot bands of high purity iron containing 0.002 wt. % C, trace Mn, 0.002 wt. % P, 0.006 wt. % S and 0.0026 wt. % O, and Fe-0.02 wt. % C alloy containing 0.16 wt. % Mn, 0.002 wt. % P, 0.0008 wt. % S. The former was used in the as hot rolled condition, while the latter w a s annealed at 700°C for 10 min, and quenched into iced brine, to keep all carbon atoms in solid solution. Both specimens were cold rolled to various rolling reductions up to 70 % in thickness under good lubrication. Thin foils prepared from the longitudinal sections of these specimens were observed with a J E O L 200 C X transmission electron microscope operated at 200 kV. RESULTS AND DISCUSSION Development of dislocation substructures during cold rolling Dislocation substructures developed during cold rolling have been intensively studied in fee metals and, at present, it is well established that microstructures observed at rolling reductions between 10 and 80 % are characterized by the formation of isolated microbands at low rolling reductions, and their clustering at high rolling r e d u c t i o n s . 8
Similar microband structures have been frequently observed also in cold rolled iron and low carbon steels ' " . However, although it is thus generally accepted that these microband structures are formed by the formation of isolated microbands and their clustering, it is not k n o w n h o w an isolated microband is nucleated within a grain. In order to form well-developed microband clusters at high rolling reductions, n u m e r o u s isolated microbands must be formed at low rolling reductions. However, the observed isolated microbands w e r e not so n u m e r o u s . Also the detailed mechanism of their clustering is not k n o w n at present. Therefore, it was strongly suggested that these microband structures were formed by a different mechanism. 9
In Figs. 1 to 5, microstructures observed with T E M on the longitudinal sections of the specimens cold rolled to various rolling reductions are illustrated. Equi-axed cells w e r e observed already in the specimen cold rolled to 10 % reduction in thickness. Figure 1 shows such cell structures formed in a {111} < 1 1 2 > deformed grain. After cold rolling to 20 % reduction in thickness, Fig. 2, these cells began to be elongated in the rolling direction. Those parts of the square cell walls lying m o r e remote from the rolling direction began to be disintegrated, while those parts of cell walls lying nearer to the rolling direction developed into long and dense dislocation walls. In this figure, it should be noted that, although elongated cells were formed in the region A, cells were still equi-axed in the region B . Thus, this process of the cell elongation seems to be propagated gradually into surrounding regions. (In this process,
194
Materials Processing and Texture
I m p o r t a n c e of D i s l o c a t i o n S u b s t r u c t u r e s in T e x t u r e C o n t r o l of L o w C a r b o n S t e e l
Fig. 1 Equi-axed cells observed on the longitudinal section of high purity iron cold rolled 10 % reduction in thickness
Fig. 2
Decomposition of equi-axed cells and their rearrangement into elongated cells. H i g h purity iron cold rolled to 20 % reduction in thickness. Longitudinal section.
disintegration of equi-axed cells is very important. Very long microbands cannot be formed by the simple elongation of equi-axed cells without their prior decomposition.) Through this continuous rearrangement of equi-axed cells into elongated cells, so called microband structures consisting of elongated cells were well developed after cold rolling of 50 % reduction in thickness. Figure 3 s h o w s such microband structures developed in a {111} <112> deformed grain.
Materials Processing and Texture
195
I m p o r t a n c e of D i s l o c a t i o n S u b s t r u c t u r e s in T e x t u r e C o n t r o l of L o w C a r b o n S t e e l
Fig. 3 Microband structures observed in (111) [112] deformed grain after cold rolling of 50 % reduction in thickness. Longitudinal section.
Fig. 4 Dislocation tangles observed in (001) [110] deformed grain after cold rolling of 50 % reduction in thickness. Longitudinal section.
196
·
Materials Processing a n d Texture
I m p o r t a n c e of D i s l o c a t i o n S u b s t r u c t u r e s in T e x t u r e C o n t r o l of L o w C a r b o n S t e e l
Thus, in this high purity i r o n , neither formation of isolated microbands nor their clustering could be observed. It is k n o w n that the sizes and morphologies of these microbands are strongly dependent on the orientation of the deformed g r a i n ' ' ~ ' \ In low carbon steels cold rolled to 70 % reduction in thickness. Every and H a t h e r l y found that equi-axed cells were formed within grains having their normal direction, N D , within 10° from < 1 0 0 > . Elongated cells were observed within grains having other orientations. Österle et a l . found that elongated microband structures were developed in deformed grains belonging to the <111 > // N D fiber textures. In deformed grains having orientations belonging to < 1 1 0 > // R D (rolling direction) fiber textures, equi-axed cells were observed. 5
9
1 0
1
9
12
In agreement with this, d e n s e dislocation tangles w e r e observed in the {001} < 1 1 0 > deformed grain after cold rolling of 50 % in thickness. Fig. 4. In this figure, it is also shown that a few long and straight dislocation walls were developed within this grain. Such an effect of grain orientations on the development of dislocation substructures was more clearly observed in the specimen cold rolled to 70 % reduction in thickness. Fig. 5. In Fig. 5, it is clearly s h o w n that the grain located at the center of the figure was severely extended by cold rolling. It should be noted that also microbands within this grain were severely extended or warped in wavy forms. Microbands located in the regions near grain boundaries or near grain boundary triple point evidently suffered from severer and m o r e complex deformation. Thus, microbands formed at low rolling reductions w e r e , like a deforming grain, heavily thinned and elongated at high rolling reductions. However, although boundaries between m i c r o b a n d s appeared to be low angle boundaries, observations at higher magnifications revealed that, at the rolling reduction of 70 % in thickness, they were not perfect grain boundaries yet. They were still dislocation walls connected with loose dislocation tangles w h i c h surrounded them. Therefore, it was suggested that they could a c c o m m o d a t e plastic deformation flexibly by cooperative m o v e m e n t of these dislocations or byrearrangement of dislocations within these walls. This might be the reasons why microbands were elongated and extended as a whole at high rolling reductions.
Fig. 5 Microstructures of high purity iron cold rolled 70 % reduction in thickness.
Materials Processing and Texture
•
197
I m p o r t a n c e of Dislocation S u b s t r u c t u r e s in T e x t u r e C o n t r o l of L o w C a r b o n S t e e l
Natures of these m i c r o b a n d s and their boundaries b e c a m e m o r e obvious by annealing these specimens at 520°C, Fig. 6. After annealing for 2 min, Fig. 6(a), dislocations within these microbands were significantly decreased. At the same time, boundaries of these m i c r o b a n d s became sharp. S o m e of these microbands w e r e coarsened. On further annealing, these microbands were developed into elongated subgrains containing few dislocations, Fig. 6 (b). Some of t h e m had g r o w n into spherical shapes and they were obviously in the earliest stage of the formation of recrystallization nuclei. In fact, after annealing for 8 min. recrystallized grains were formed within such groups of elongated subgrains, Fig. 6 (c).
Fig. 6 Microstructures observed o n the longitudinal sections of high purity iron cold rolled to 70 % reduction in thickness and annealed at 520 ° C for (a) 2 min (b) 4 min, and (c) 8 min.
198
Materials Processing a n d Texture
I m p o r t a n c e of Dislocation S u b s t r u c t u r e s in T e x t u r e C o n t r o l of L o w C a r b o n S t e e l
These results clearly s h o w that microbands are very important microstructure elements providing recrystallization nucle. Control of these microband structures is thus the key technology in texture control. In the specimen cold rolled to 70 % reduction in thickness, another type of bands was often observed, Fig. 7. They w e r e very long and thin, m a k i n g an angle of 34 ° to the rolling direction (see arrows). These may be shear bands which have been frequently observed in other i n v e s t i g a t i o n s " ' . In fact, at these points where they cross microbands, large shear deformation w a s a l w a y s introduced into the latter. 9
14
lo
Fig. 7 Microstructures observed on the longitudinal section of high purity iron cold rolled to 70 % reduction in thickness. A r r o w s show shear bands. T h e development of such dislocation substructures should be strongly influenced by alloying elements and second phase particles. The effect should be strongest in the case of the carbon atom in solid solution. This is because in iron and steels, interstitial carbon atoms interact most strongly with dislocations, inducing frequent cross slip during deformation. In previous i n v e s t i g a t i o n s ' , this effect of the solute carbon a t o m was investigated in detail in 0.02 and 0.04 wt. % C steels solution treated at 700°C for 10 min. quenched into iced brine, and cold rolled up to 90 % reduction in thickness. T E M observations m a d e on these specimens revealed that, at low rolling reductions, solute carbon atom suppressed the formation of normal equi-axed cell structures. Thus, only ill-defined equi-axed cells with loose and broad cell walls w e r e observed. On further cold rolling, these equi-axed cells were rearranged into dense dislocation tangles elongated in one direction. After cold rolling of 70 % reduction in thickness, all area w a s covered with dense dislocation tangles. The development of microband structures w a s thus strongly suppressed by the solute carbon atom. 7
, 8
In this paper, this effect of the solute carbon atom was investigated more in detail. Figure 8 illustrates microstructures observed on the longitudinal section of Fe-0.02 wt. % C alloycontaining an excess amount of carbon atoms, and cold rolled to a reduction of 70 % in thickness.
Materials Processing a n d Texture
199
I m p o r t a n c e of Dislocation S u b s t r u c t u r e s in T e x t u r e C o n t r o l of L o w C a r b o n S t e e l
It is evident that, also in this specimen, the development of microband structures w a s strongly suppressed, and dense dislocation tangles were distributed h o m o g e n e o u s l y . Only illdefined microbands could be observed locally. Microstructures of iron and steels containing an excessive a m o u n t of carbon atoms in solid solution and cold rolled to a reduction of 70 % in thickness were characterized by the formation of several types of banded s t r u c t u r e s ' ' . In Fig. 8, it is seen that very thin, straight and long bands (see arrows) were nucleated at the grain boundary ( C B ) . They propagated into the grain interior, running parallel to each other. Figure 9 illustrates shear bands (see arrows) observed in another area of the same specimen. They were nucleated at the neck (see arrow) formed on the grain boundary ( G B ) , and g r e w into the grain interior. They w e r e curved, showing frequent branching. Most of them w e r e much broader than the shear bands given in Fig. 5. Carbon atoms in solid solution thus enhanced the formation of such broad shear bands. 7
8
Fig. 8 Dislocation substructures formed in Fe-0.02 wt. % C alloy solution treated at 700 ° C for 10 min. quenched into iced brine, and cold rolled to 70 % reduction in thickness. A r r o w s show thin bands. G B : Grain boundary. These t w o types of b a n d s are s h o w n in Fig. 10 in enlarged scales. It is seen that a thin band consisted of a single, long band, whereas a shear band c o m p o s e d of a bundle of elongated subgrains. which were m u c h shorter than the thin band. These b a n d s could be observed also with an optical microscope. Figure 11 (a) illustrates these bands observed on the longitudinal section of the cold rolled specimen. A broad band traversing the deformed grain and m a k i n g an angle of about 30° to the rolling direction (horizontal) is clearly a shear band. T w o types of faint and thin strain markings m a y b e thin bands observed with T E M . O n e w a s almost parallel to the broad shear band, while the other m a d e angle of about 20° to the rolling direction. Both were interlaced with each odier. yielding herring bone patterns.
200
Materials Processing a n d Texture
I m p o r t a n c e of Dislocation S u b s t r u c t u r e s in T e x t u r e C o n t r o l of L o w C a r b o n S t e e l
Fig. 9 Shear b a n d s SB (white arrows) nucleated at the neck o n the grain boundary, G B . Fe-0.02 wt. % C alloy solution treated at 700 ° C for 10 m i n , quenched into iced brine, and cold rolled to 70 % reduction in thickness.
Fig. 10
(a) Thin bands and (b) shear band observed at high magnifications.
Materials Processing and Texture
·
201
I m p o r t a n c e of D i s l o c a t i o n S u b s t r u c t u r e s in T e x t u r e C o n t r o l of L o w C a r b o n S t e e l
Fig. 11 (a) Shear b a n d s and thin b a n d s observed with a n optical microscope. Fe-0.02 wt. % C alloy solution treated at 700 ° C for 10 min, quenched into iced brine, and cold rolled to 70 % reduction in thickness, (b) Nucleation of recystallized grains at shear bands observed after annealing at 520 °C for 4 min. 16
In high purity iron cold rolled to 80 % reduction in thickness, N i n g et a l . already observed two types of bands, which w e r e inclined 17° and 32° to the rolling direction. Thus, the observed results agreed quite well with those reported by N i n g et al. On annealing the cold rolled specimen at 520 °C for 2 m i n , recrystallized grains w e r e formed only at broad shear bands. Figs. 11 (b) and 12. T h e width of a single thin band is clearly smaller than the critical radius of the recrystallization nuclei, whereas a broad shear band was wide enough to be a preferred nucleation site of a recrystallized grain. Formation of these bands depended strongly on the orientation of the deformed grain, Fig. 13. It is seen that broad shear bands which were nucleated at the gtain boundary (see arrow), traversed several deformed grains, inducing offsets at their grain boundaries. In the deformed grain in which broad shear bands were nucleated, also thin bands could be observed. They were also nucleated at the grain boundary. O n the other hand, some deformed grains were completely free from such bands. The morphologies of broad shear b a n d s illustrated in Fig. 13 were very similar to those observed in an Al-5 wt. % M g alloy cold rolled to 50 % reduction in t h i c k n e s s " . Contribution of dislocation substructures to the development of cold rolling textures In the past, the development of cold rolling textures in metals and alloys has been interpreted in most cases in terms of T a y l o r ' s h o m o g e n e o u s deformation t h e o r y . In fee metals such as Al, Cu and Brass, the rotation paths and stable end orientations have been adequately predicted with this t h e o r y . 20
21
202
Materials Processing a n d Texture
I m p o r t a n c e of Dislocation S u b s t r u c t u r e s in T e x t u r e C o n t r o l of L o w C a r b o n S t e e l
Fig. 12 Nucleation of recystallized grains at a shear band (SB) o b s e r v e d with T E M .
Fig. 13 Shear bands and thin bands formed in high purity iron polycrystals.(Arrow shows the nucleation site of a broad shearband.) 2 2
2 6
A l s o in iron and low carbon steels, several attempts have been m a d e to simulate the development of cold rolling textures with T a y l o r ' s theory. However, since these theoretical analyses are based o n the h o m o g e n e o u s , macroscopic plasticity theory, a microscopic process such as the development of dislocation substructures during cold rolling has not been taken into account. The development of the cold rolling texture may be influenced by the formation of dislocation substructures, since the shape and dimensions of microbands prescribe the mean free path of m o v i n g dislocations inducing slip rotation. These microbands have an anisotropic, pancake shape, elongated in the rolling direction. Therefore, the m e a n free path of dislocations should be determined by the shortest dimension, i. e., the thickness of the microband. With increasing rolling reduction, the thickness of the microband, i. e. the m e a n free path of dislocation decreases, so that the development of cold
Materials Processing and Texture
203
I m p o r t a n c e of D i s l o c a t i o n S u b s t r u c t u r e s in T e x t u r e C o n t r o l of L o w C a r b o n S t e e l
rolling textures becomes gradual. T h u s , except at high rolling reductions, where d y n a m i c recovery occurs appreciably, the thickness of the microband might be the most important parameter that controls the development of cold rolling textures. Rolling textures are d e v e l o p e d d u e t o slip rotation, w h i c h o c c u r s during cold rolling. Such rotation occurs along a definite path determined by active slip systems and initial orientations of crystals. In previous i n v e s t i g a t i o n s ~ , such rotation paths were determined by using iron polycrystalline specimens having various strong initial textures. By cold rolling these specimens to various rolling reduction, and by determining main orientations at each rolling reduction with the orientation distribution function analysis, it w a s found that, during cold rolling, individual grains in polycrystals were rotated along one of the following two paths; (A) {110} <001> {111} < 1 1 2 > {111} < 1 1 0 > { 2 2 3 } < 1 1 0 > (B) {001} < 1 0 0 > {001} < 1 1 0 > {112} < 1 1 0 > { 2 2 3 } < 1 1 0 > aiming at c o m m o n {223} <110> stable end orientations. 27
29
It was further found that the rotation along the path A occurred very rapidly during cold rolling. Grains h a v i n g {110} <001 > orientations could arrive at {111} < 1 1 0 > already after cold rolling of 70 % reduction in thickness, and at {223} < 1 1 0 > stable end orientations after cold rolling of 90 % reduction in thickness. On the other hand, grains having {001} <110> orientations rotated very slowly along the path Β during cold rolling. This is because dense dislocation tangles were developed h o m o g e n e o u s l y within such grains, Fig. 2. Thus, during cold rolling, most of texture changes occurred by the rotation along the path A. Since m i c r o b a n d s w e r e formed in the grain rotating a l o n g the path A, the development of t h e cold rolling texture should be controlled for the most part by the development of the microband structures. In the past, it has been well established that, during cold rolling, the development of the micoband structures occurs extensively at rolling reduction above 20 % in thickness . On the other hand, it has been also well recognized that the development of cold rolling textures occurs most significantly in the s a m e r a n g e o f the rolling reduction. Since these t w o processes occur in the same specimen, it is quite reasonable to conclude that the development of cold rolling textures is controlled by the development of such microband structures. 8
In high purity iron, it has been observed that the intensity of the cold rolling texture can be varied as m u c h as a factor 2 by the addition of the alloying elements ' . Such an effect of the alloying addition seems to be derived from the fact that the alloying element affects the development of microband structutures . 3 0
3 I
7
The most critical step in the formation of such microband structure is the disintegration of equi-axed cells and their rearrangement into elongated cells, Fig. 2. This process has been already observed by Ahlhorn and S a u e r in Cu single crystals subjected to wire drawing. 32
CONCLUSIONS (1) So called microband structures are formed during cold rolling by the disintegration of equiaxed cells and their subsequent rearrangement into elongated cell structures. At high rolling reductions, these elongated cells are, like, deforming grains, heavily deformed and extended during cold rolling.
204
Materials Processing a n d Texture
I m p o r t a n c e of Dislocation S u b s t r u c t u r e s in T e x t u r e C o n t r o l of L o w C a r b o n S t e e l
(2) Interstitial carbon atoms significantly suppress the development of such microband structures. (3) The development of cold rolling textures is closely related with the development of dislocation substructures. This is because the shortest dimension of the microband. i. e. its thickness determines the mean free path of dislocations inducing slip rotation. Since the thickness of microbands decreases with increasing rolling reduction, slip rotation becomes m o r e difficult at high rolling reductions. A l s o d y n a m i c recovery suppresses slip rotation at high rolling reductions. A s a result, the texture development b e c o m e s gradual at high rolling reductions.
REFERENCES Ή. Inagaki, K. Kurihara, and I. Kozasu, ' I n f l u e n c e of Crystallograhic Texture on the Strength and T o u g h n e s s of Control-Rolled H i g h Strength Steel," Trans. ISU, 17, 75-81 (1977). H . Inagaki, "Effect of Crystallographic Texture on the Separration Behavior of ControlRolled L o w Carbon Steel," Z. Melallkd.. 7 6 . 65-90 (1985). Ή . Inagaki. " Fractographic Study of the Formation of the Separation in a Control-Rolled L o w Carbon N b Steel," Ζ Metallkd., 7 9 , 364-374 (1988). H . Inagaki. "Effect of the Crystallographic Texture on the Ultrasonic Propagation in Control-Rolled L o w Carbon Steels." Ζ Metallkd.. 8 0 , 4 1 6 - 4 2 4 (1989). ^W. B. Hutchinson, " M e c h a n i s m of Recrystallization and Texture Evolution in Cold Rolled Steels," Proc. of 21" Riso Symposium on Materials Science, ed. by N . Hansen, X. Huang, D. Juul Jensen, Ε. M . Laudrisen, T. Leffers, P. Pantelon, T. J. Sabin, and J. A. Wert, Riso Nat. Lab., Roskilde. D e n m a r k , 91-102 (2000). 2
4
Ä
R . K. R a y . J. J. Jonas, and R. E. Hook, "Cold Rolling and Annealing Textures in L o w Carbon and Extra L o w Carbon Steels." Int. Mater. Review. 3 9 . 129-172 ( 1977). H . Inagaki, "Fundamental Aspect of Texture Formation in L o w Carbon Steel," ISU Int., 3 4 . 3 1 3 - 3 2 1 (1994). A . S. Malin, and M . Hatherly, "Microstructures of Cold-Rolled Copper. " Metal Science, 13,463-472(1979). R . L. Every, and M . Hatherly, "Oriented Nucleation in L o w Carbon Steels. " Texture. 1. 183-194(1974). D . J. Willis, and M . Hatherly, "Orientation D e p e n d e n c e of Cell Sizes in N b Stabilized Steels and the Development of the Recrystallizatioon Texture, Proc. j " ' ICOTOM, ed. by G. Gottstein and K. Lücke, Springer Verlag, Berlin. 465-474 (1978). ' 'R. L. Agahn. and J. Nutting, "Structure and Properties of Low Carbon Steels after Deformation to High Strains," Mtal Science, 14, 233-237 (1980). W. Österle, H. Wever, and H. J. B u n g e , "Development of Microstructures and Texture of Cold Rolled α Iron," Metal Science. 1 7 . 333-340 (1983). ° H . J. Roven, and E. N e s . "Structure and Properties of High Strength L o w Alloy and L o w C a r b o n Steels, " Metal Sei., 1 8 . 5 1 5 - 5 2 0 ( 1 9 8 4 ) . M . Feller-Kniepmeier, and N . Wanderka, "Elektronenmikroskopische Untersuchung der Bildung von Mikrobändern und Scherbändern in Gewalzten Nickel und α Eisenkristallen, " Ζ Melallkd, 7 9 , 673-677 (1988). 7
8
9
I 0
1 2
I 4
Materials Processing and Texture
·
205
I m p o r t a n c e of Dislocation S u b s t r u c t u r e s in T e x t u r e C o n t r o l of L o w C a r b o n S t e e l
' \ ) . Liu. B. L. Li, W. Liu. and X. H u a n g , "Deformation Structure Evolution of IF Steel during Cold-Rolling, " Reciyslallizalion-Fundamenlal Aspects and Relations to Deformation Microstructure, ed. by N . H a n s e n , X. H u a n g , D. Juul Jensen, E. M . Laudrisen, T. Leiters, P. Pantelon. T. J. Sabin, and J. A. Wert. Riso Nat. Lab.. Roskilde. Denmark, 423-430 (2000). I 6
-
H . Ning, B . J. Duggan, G. L.Liu, and X. Z h a n g . ' The Effect of Shear B a n d s on the Evolution of Recrystallization Textures in Pure Iron." Prot: ICOTOM 11. ed. by Υ Ζ. Liang, L. Zou, and Y. Chu, Int. Acad. Publ.. Beijin. 525-530 (1990). H . Inagaki, "Effect of Dissolved Carbon Atom on the Development of the Rolling Texture in Polycrystalline Iron." Z. Metallkd, 8 1 . 15-24 (1990). H . Inagaki. "Development of Microstructures and Textures During Cold Rolling of Polycrystalline Iron Containing an Excess A m o u n t of Dissolved Carbon." Z. Metallkd, 8 1 , 4 7 4 483(1990). H . Inagaki, and S. K.ohara, "Shear bands in Cold Rolled A l - M g Alloy Polycrystalst." Z. Metallkd. 8 8 , 570-575 ( 1997). G . I. Taylor, "Plasic Strain in M e t a l s , " J. Inst. Metals. 62, 307-324 (1938). J . Hirsch, and K.. Lücke, "Mechanism of Deformation and D e v e l o p m e n t of Rolling Textures in Polycrystalline F C C Metals-II Simulation and Interpretation of Experiments on the Basis of Taylor-Type Theories," Acta metall, 3 6 , 2883-2904 (1988). I . L. Dillamore and H. Katoh, "Polycrystalline Plasticity and Texture Development in Cubic Metals. " BISRA Report. British Steel Corp., London, M G / 1 3 / 7 1 , 1-17(1972). T . L. Raphanel, and P. V a n Houtte, "Simulation of the Rolling Textures of B C C Metals by M e a n s of the Relaxed Taylor Theory, "Acta metal!., 3 3 . 1481-1488 (1985). P. V a n Houtte and L. Rabet, '"Generalization of the relaxed constraints models for the prediction of deformation textures." Revue de Metallurgie-Cathiers D Informations Techniques, 94,1483-1494(1997). P. Van Houtte, S. Y. L i . M . Seefeldt and L. Delannay,"Deformation texture prediction : from the Taylor model to the advanced Lamel m o d e l , " Int. J. Plasticity 2 1 , 589-624 (2005). L . S. Toth, J. J. Jonas, D. Daniel, and R. K. Ray, "Development of Ferrite Rolling Textures in Low and Extra L o w Carbon Steels. " Metall. Trans.. 21A, 2958-3000 (1988). H . Inagaki. and T. Suda, "The Development of Rolling Textures in L o w Carbon Steels" Texture, 1, 1 2 9 - 1 4 0 ( 1 9 7 2 ) . H . Inagaki. "Formation of {111} Recrystallization Textures in Polycrystallline Fe," Trans. ISIJ. 2 4 . 266-274 (1984). H . Inagaki. "Stable End Orientations in the Cold Rolling Textures of the Polycrystalline Iron." Z. Metallkd, 7 8 . 431 - 4 3 9 ( 1987). " H. Inagaki, "Effect of the Alloying Element on the Development of the Rolling Texture in High Purity Iron." Ζ Metallkd, 7 9 , 164-173 (1988). H . Inagaki, "Effect of Ti on the Development of Rolling Textures in High Purity Iron," Textures and Microstr., 8 & 9 . 173-189 (1988). H . Ahlhorn and D. sauer. "Elektronen-mikroskopische Untersuchungen gezogener Kupfereinkristalle." Ζ Metallkd, 5 9 , 658-663 (1968). 1 7
1 8
1 9
2 0
2 1
2 2
2 ,
2 4
2 5
2 o
2 7
2 8
2 ,
1fl
3 I
1 2
206
Materials Processing and Texture
T E X T U R E C O N T R O L IN M A N U F A C T U R I N G C U R R E N T A N D F U T U R E G R A D E S O F LOW-CARBON STEEL SHEET Leo A.I. K e s t e n s ' and R o u m e n Petrov'
' Materials Science and Engineering Department Delft University of Technology Mekelweg 2 2628 C D Delft, The Netherlands 2
Department of Metallurgy and Materials Science Ghent University (UGent) Technologiepark 903 B-9052 Ghent, Belgium
ABSTRACT Quantitative texture analysis has significantly contributed to the development of new steel qualities in the last decades. For a n u m b e r of steel grades (e.g. deep drawing and electrical steels) the relevant properties are explicitly dependent on the crystallographic anisotropy. It is self-evident that for these steels the precise control of texture development in the subsequent stages of the manufacturing process is of crucial importance. But also for the products of which the properties do not bear a close connection with the crystallographic texture, the assessment of textural change in solid-state transformation may contribute to a better understanding of the formation of microstructures and thus an improved control of all structure sensitive properties. In this paper an overview is presented of the current state-of-the-art and future trends with regard to texture control in deep drawing steels and advanced high strength steels such as T R I P and T W I P steels. T h e potential of the double cold rolling process with intermediate annealing is discussed for the development of super deep drawable grades. It is shown that a selective transformation of {111 }<110> orientations from the recrystallized ferrite matrix m a y occur during the intercritical annealing of a cold rolled T R I P steel. The recrystallization process in cold rolled T W I P steels is largely conditioned by the stacking fault energy of the material and specific features of the deformation structure. KEYWORDS Texture control, low-carbon steels, deepdrawability, T R I P steel, T W I P steel.
207
T e x t u r e C o n t r o l in M a n u f a c t u r i n g C u r r e n t a n d F u t u r e G r a d e s of L o w - C a r b o n S t e e l S h e e t
INTRODUCTION Since the seventies quantitative texture analysis m e t h o d s have been applied in a wide variety of materials science applications. O n e of the main achievements o f the texture research c o m m u n i t y is that n e w metals, and in particular n e w steel alloys, have been developed with tailored textures so as to optimize their mechanical and functional properties. O n a m o r e fundamental level the detailed analysis of texture formation m e c h a n i s m s has provided a unique and revealing entry to the field of physical metallurgy. The investigation of the crystallographic features of various solid-state transformation processes has led to a substantial progress in our understanding o f the microstructural transformation m e c h a n i s m s . This is true for p h e n o m e n a related to phase transformation , plastic deformation and ensuing softening processes such as recrystallization and grain growth · . The aim of the present paper is not so m u c h to present a crystallographically resolved description of the materials science physics involved in these phenomena, but rather to describe a n u m b e r of e x a m p l e s in which the texture plays an instrumental role in developing n e w steel products. 1
2
S T E E L S FOR D E E P D R A W I N G A P P L I C A T I O N S D e e p drawing is a particular form of sheet forming which can be modelled in a laboratory experiment by a cup drawing test. The deepdrawability of a metal sheet is said to improve w h e n circular blanks with increasing diameters can be drawn to cups without failure. For single p h a s e materials the deepdrawability has a close relation with the normal anisotropy r which represents the ratio of the width-to-thickness strain in a tensile test. By analysing the shape of the yield locus one can derive that the deepdrawability benefits from a large normal anisotropy. Conversely, the in-plane anisotropy (i.e. the Ar = (ΓΟ+Γ9Ο-2Γ )/2) should be m i n i m u m for o p t i m u m metal forming operations. It w a s shown both empirically and theoretically that the normal and planar anisotropy critically depend on the texture. For B C C materials the double boundary condition of m a x i m i z i n g the average normal anisotropy r ( r = (ro+2r +r9o)/4) and minimizing the planar anisotropy Ar can be complied with by an ideal γ-fibre texture in which all orientations exhibit a {111} crystal plane parallel to the rolling plane because this texture gives rise to an appropriate yield locus with optimized shape for deepdrawing. 45
m
m
45
In the seventies it w a s discovered that strong γ-fibre recrystallization textures could be produced by severely reducing the content of carbon interstitial solutes in the Fe lattice. In interstitial free (IF) steels the latter is c o m m o n l y obtained by v a c u u m degassing the converter load and alloying the steel with small additions of Ti and/or 