Additive and Traditionally Manufactured Components: A Comparative Analysis of Mechanical Properties (Additive Manufacturing Materials and Technologies) [1 ed.] 0128219181, 9780128219188

Table of contents :
Cover
ADDITIVE AND
TRADITIONALLY
MANUFACTURED
COMPONENTS
A Comparative Analysis of Mechanical
Properties
Copyright
Dedication
Preface
About the author
1
What is additive manufacturing?
2
Fabrication
Fused deposition method (FDM)
Melt properties
Liquefier
Heat convection
Pressure drop estimation
Layer deposition and stability
Road spreading
Road cooling and polymer bonding
Powder-bed fusion (PBF)
Inkjet printing
Stereolithography (SLA)
The state of the resin (photopolymer)
The maximum cure depth
The cured line width
Laser scan velocity
Direct energy deposition (DED)
Thermal model
Laminated object manufacturing (LOM)
References
Further reading
3
Testing: Comparison of AM data with traditionally fabricated
Tensile tests
Ti-6Al-4V: AM tensile properties
Al alloy AA6061: AM tensile properties
Conventionally produced (AM) AA6061
Stainless steel 304L: AM tensile properties
Conventionally produced SS 304L
Ceramic
AM alumina
Conventionally fabricated alumina
Compression tests
Ti-6Al-4V
Conventionally fabricated Ti-6Al-4V
Al alloys-Al 60613
Conventionally fabricated Al 6061
AM stainless steel 304L
Conventionally fabricated stainless steel 304L
Ceramics-Alumina
Conventionally fabricated alumina (Al2O3)
Effect of orientation and temperature
Indentation (hardness)
Ti-6Al-4V
Conventionally produced Ti-6Al-4V
Aluminum alloy (Al6061)
Conventionally fabricated Al 6061
Stainless steel 304L
Conventionally produced 304L stainless steel
Alumina
Conventionally produced alumina
Temperature dependence
Hardness of coatings
Hardness of alumina films
References
Further reading
4
Dislocations in AM and traditional manufacturing: A comparison
Introduction
In AM Ti-6Al-4V
In traditionally fabricated Ti-6Al-4V
Motion of dislocations
Introduction AA6061
AM of AA6061 Al alloy
Dislocations in conventionally produced Al AA6061
Pinning of dislocations in 6061
The strain effect in 6061
In stainless steel 304L
Introduction
In AM 304L stainless steel
In conventionally fabricated 304L stainless steel
In alumina (Al2O3)
In conventionally fabricated alumina
References
Further reading
5
Deformation in AM and traditional manufacturing: A comparison
Introduction
Deformation in AM Ti-6Al-4V
In traditionally fabricated Ti-6Al-4V
Tensile deformation
Compressive deformation
Deformation in AM Al AA6061
Tensile deformation in Al AA6061
Compressive deformation
Conventional tensile deformation
Conventional compressive deformation
AM stainless steel 304L
Tensile deformation
Compression deformation
Conventionally produced SS 304L
Tensile deformation in conventionally produced SS 304L
Compressive deformation in conventionally produced SS 304L
Deformation in alumina
Compressive deformation of AM alumina
Hardness
References
Further reading
6
Dynamic deformation
Introduction
Dynamic deformation of AM Ti-6Al-4V
Tensile test of AM Ti-6Al-4V
Tensile test of CP Ti-6Al-4V
Compression tests
In AM Ti-6Al-4V
In CP Ti-6Al-4V
Twinning in Ti-6Al-4V
Dynamic deformation in Al AA6061
Tension test in AM AlSi10Mg
Compression test in AM Al Si10Mg
Tensile test in CP AA6061
Compression test in CP Al 6061
Tensile test in AM SS 304L
Compression test in AM SS 304L
Tensile test in CP 304L SS
Compression test in CP 304L SS
Dynamic deformation in alumina (Al2O3)
Tension test in AM alumina
Compression test in AM alumina
Hardness in AM alumina
Tensile test in CP alumina (Al2O3)
Compression test in CP alumina (Al2O3)
References
Further reading
7
Time-dependent deformation creep in AM and traditional manufacturing
Introduction
Tensile creep in AM Ti6Al4V
Compressive creep in AM Ti6Al4V
Tensile creep in CP Ti6Al4V
Compressive creep in CP Ti6Al4V
Tensile creep in AM Al10SiMg
Tensile creep in CP Al AA6061
Compressive creep in CP Al AA6061
References
Further reading
8
Cyclic deformation (fatigue) in AM and traditional manufacturing: A comparison
Introduction to fatigue
Fatigue in AM Ti6Al4V
High cycle fatigue
Low cycle fatigue
Rough surface and notch effect
Fatigue in conventionally fabricated Ti6Al4V
High cycle fatigue
Low cycle fatigue
Rough surface and notch effect
Fatigue in conventionally fabricated Al AA6061
High cycle fatigue in Al 6061
Low cycle fatigue
The Massing hypothesis
Rough surface and notch effect
Fatigue in AM SS 304L
Hgh cycle fatigue
Fatigue in CP SS 304L
High cycle fatigue
References
Further Reading
9
Fracture in AM and traditional manufactured components
Fracture in AM Ti-6Al-4V
Fracture in AM Al AA6061
Fracture in AM SS 316L
Fracture in AM alumina
Fracture in CP Ti-6Al-4V
Fracture in CP Al AA6061
Fracture in CP SS 304L
Strain rate effects in CP SS 304L
Hydrogen effects in CP SS 304L-Hydrogen embrittlement
Introduction
Fracture in CP alumina
References
Further reading
10
Comparison of deformation in AM and CP nanomaterials
Tensile properties
AM Ti6Al4V
CP Ti6Al4V
AM of nano-316L SS
CP nano-316L SS
CP nano-316L and 304L SS
CP nano-304L SS
Compressive properties
AM of nano-alumina
CP of nano-alumina
Indentation hardness in nanomaterials
Introduction
Hardness in AM nano-alumina
Hardness in CP nano-alumina
References
Further reading
Epilogue
Index
A
B
C
D
E
F
G
H
I
L
M
N
P
S
T
V
Back Cover

Citation preview

ADDITIVE AND TRADITIONALLY MANUFACTURED COMPONENTS A Comparative Analysis of Mechanical Properties

ADDITIVE AND TRADITIONALLY MANUFACTURED COMPONENTS A Comparative Analysis of Mechanical Properties

JOSHUA PELLEG

Materials Engineering Department Ben Gurion University of the Negev, Beer Sheva, Israel

Elsevier Radarweg 29, PO Box 211, 1000 AE Amsterdam, Netherlands The Boulevard, Langford Lane, Kidlington, Oxford OX5 1GB, United Kingdom 50 Hampshire Street, 5th Floor, Cambridge, MA 02139, United States © 2020 Elsevier Inc. All rights reserved. No part of this publication may be reproduced or transmitted in any form or by any means, electronic or mechanical, including photocopying, recording, or any information storage and retrieval system, without permission in writing from the publisher. Details on how to seek permission, further information about the Publisher’s permissions policies and our arrangements with organizations such as the Copyright Clearance Center and the Copyright Licensing Agency, can be found at our website: www.elsevier.com/permissions. This book and the individual contributions contained in it are protected under copyright by the Publisher (other than as may be noted herein). Notices Knowledge and best practice in this field are constantly changing. As new research and experience broaden our understanding, changes in research methods, professional practices, or medical treatment may become necessary. Practitioners and researchers must always rely on their own experience and knowledge in evaluating and using any information, methods, compounds, or experiments described herein. In using such information or methods they should be mindful of their own safety and the safety of others, including parties for whom they have a professional responsibility. To the fullest extent of the law, neither the Publisher nor the authors, contributors, or editors, assume any liability for any injury and/or damage to persons or property as a matter of products liability, negligence or otherwise, or from any use or operation of any methods, products, instructions, or ideas contained in the material herein. Library of Congress Cataloging-in-Publication Data A catalog record for this book is available from the Library of Congress British Library Cataloguing-in-Publication Data A catalogue record for this book is available from the British Library ISBN: 978-0-12-821918-8 For information on all Elsevier publications visit our website at https://www.elsevier.com/books-and-journals

Publisher: Matthew Deans Acquisitions Editor: Christina Gifford Editorial Project Manager: Ana Claudia A. Garcia Production Project Manager: Poulouse Joseph Cover designer: Greg Harris Typeset by SPi Global, India

Dedication To my wife Ada and my children, Deenah and her late husband Gidon Barak, Ruth and Christer Kallevag, Shlomit and Asher Pelleg and to my grandchildren: Roy, Tal, Rotem and Noa Barak; Ella and Maya Kallevag; and Ofir and Ori Pelleg.

Preface The purpose of this book is to test the mechanical properties of materials produced by additive manufacturing (AM) and compare the test results with test values obtained by conventionally produced (CP) methods. Since the production of objects by the AM technique is for various materials, from paper through plastics to metallic alloys and ceramics, samples selected for the comparison cover the same materials, namely, polymers, alloys, and ceramics. The alloys chosen are those which are commonly known to be produced by AM. Thus the alloys considered are Al alloys, Ti alloys, and steel. The chapters considered in this book are outlined in the content. Accordingly, Chapter 1 defines shortly the AM process listing the six main type techniques. These are: fused deposition method (FDM), powder-bed fusion (PBF), inkjet printing, stereolithography (SLA), direct energy deposition (DED), and laminated object manufacturing (LOM). A detailed description of these techniques is presented in Chapter 2 dealing with fabrication. Each section, representing the above AM processes is accompanied with mathematical expressions, which can be utilized for experimental design. In Chapter 3, testing by tension, compression and indentation (hardness) of the AM fabricated specimens is considered. The test results are compared with similar tests performed on specimens obtained by conventional fabrication methods. Deformation phenomena in specimens fabricated by AM are discussed in Chapter 5. This follows the basic concept of dislocations which are responsible for the deformation in the samples. It sets the theoretical background of the experimental observations. Following the basic concept related to deformation and responsible for their occurrence under stress, namely dislocations, deformation phenomena in specimens fabricated by AM are discussed in Chapter 5. Static deformation, time-dependent deformation (creep), cyclic deformation (fatigue), and fracture in specimens obtained by AM fabrication are presented in Chapters 6–9, respectively. Again similar test results obtained by conventional fabrication are compared with those of the AM specimens. The mechanical properties of nanoscale specimens obtained by AM and traditional fabrication are presented in Chapter 10. The book closes with Chapter 11 an epilogue summarizing the observations in AM and CP materials.

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Preface

I would like to express my gratitude to all publishers and authors for permission to use and reproduce some of their illustrations and microstructures. Finally, without the tireless devotion, help, understanding, and unlimited patience of my wife Ada, I could never have completed this book, despite my decades of teaching in this field; her encouragement was essential and her helpful attitude was instrumental in inspiring to write this book.

About the author Joshua Pelleg received his BSc in chemical engineering from the Technion— Institute of Technology, Haifa, Israel; an MSc in metallurgy from the Illinois Institute of Technology, Chicago, IL and a PhD in metallurgy from the University of Wisconsin, Madison, WI. He has been in the Materials Engineering Department, Ben-Gurion University of the Negev (BGU), Beer-Sheva, Israel since 1970, was among the founders of the department, and served as its second chairman. Professor Pelleg is the recipient of the Samuel Ayrton Chair in metallurgy. His specializations are in the mechanical properties of materials and the diffusion and defects in solids. He has chaired several university committees and served four terms as the chairman of Advanced Studies at Ben-Gurion University of the Negev. Prior to his work at BGU, Pelleg served as assistant professor and then as an associate professor in the Department of Materials and Metallurgy at the University of Kansas, Lawrence, KS. Professor Pelleg was also a visiting professor in the Department of Metallurgy at Iowa State University; at the Institute for Atomic Research, US Atomic Energy Commission, Ames, IA; at McGill University, Montreal, QC; in the Department of Applied Electronics, Tokyo Institute of Technology, Yokohama, Japan; and in the Department of Physics, Curtin University, Perth, Australia. His nonacademic research and industrial experience includes: chief metallurgist in Urdan Metallurgical Works Ltd., Netanyah, Israel; research engineer in International Harvester Manufacturing Research, Chicago, IL; associate research officer for the National Research Council of Canada, Structures and Materials, National Aeronautical Establishment, Ottawa, ON; physics senior research scientist, Nuclear Research Center, Beer-Sheva, Israel; Materials Science Division, Argonne National Labs, Argonne, IL; Atomic Energy of Canada, Chalk River, ON; visiting scientist, CSIR, National Accelerator Centre, Van de Graaf Group Faure, South Africa; Bell

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About the author

Laboratories, Murray Hill, NJ; and GTE Laboratories, Waltham, MA. His current research interests include mechanical properties, diffusion in solids, thin film deposition and properties (mostly by sputtering), and the characterization of thin films, among them various silicides.

CHAPTER ONE

What is additive manufacturing? Additive manufacturing (AM) also known as three-dimensional (3D) printing or 3D fabrication describes technologies to build objects by adding layer upon layer of material. Structures are made by this technique by the addition of layers, which combine to form the final product. The 3D fabrication method is universal to all material classes such as plastic, metallic, concrete, and biomaterials. The technology has been developing rapidly since its invention in the 1980s when the first technique was applied to print plastic objects. Early additive manufacturing (AM) equipment were developed at this time. One could consider truly without exaggeration that the AM technology is a manufacturing revolution characterizing the 20th century modifying all concepts and imagination of production on the industrial progress scale. The 3D printing mentioned previously is a general expression often used to indicate all types of additive manufacturing. The 3D printing is defined as a method for the fabrication of objects through the deposition of material using a print head, nozzle, or other printer technology. In the fabrication process, the computer and computer-added design (CAD) software are used to transfer the directions to the printer, which prints the objects in the desired shape and size one thin layer at a time. The layers are repeatedly printed on top of each other, being consolidated during the process until the object (shape and size) is complete. A variety of shapes can be produced in a wide range of sizes without the need of using additional operations such as drilling, extrusion, welding, etc., and thus intricate forms can be produced in one single operational step. A cycle of object processing takes a relatively short time and an important characteristic of a process is that changes can be made easily by the computer during fabrication. Such a flexibility without the need of additional mechanical operations results in considerable cost savings.

Additive and Traditionally Manufactured Components https://doi.org/10.1016/B978-0-12-821918-8.00001-2

© 2020 Elsevier Inc. All rights reserved.

1

2

Additive and traditionally manufactured components

A classification of the various technologies is presented in this chapter, and the production details of these methods are described in the chapter on fabrication (Chapter 2). The main AM methods are as follows: (1) Fused deposition method (FDM). Thermoplastic polymers are the material that are generally fabricated by this method. (2) Powder-bed fusion (PBF). It is suitable for printing complex structures. The powders in each layer are fused together with a laser beam or a binder. (3) Inkjet printing. It is one of the principal methods of ceramics manufacturing. (4) Stereolithography (STL). UV (ultraviolet) light (or electron beam) is applied to initiate a chain reaction in a resin-like or monomer solution. It is effective for complex nanocomposite production. (5) Direct energy deposition (DED). The sources of energy are laser electron beam and arc. It is used also in the automotive and aerospace industries. (6) Laminated object manufacturing (LOM). It is useful for thermal bonding of ceramics and metallic materials. In DED, energy is directed to a small region of the material to be deposited, which melts it and the requirement is that the laser energy (or other energy source) is powerful enough to melt the powder. The energy source (laser > electron beam > arc) and the deposition rate determine the quality of the DED process. Materials that can be fabricated are stainless steels, copper, aluminum and its alloys, titanium, nickel, cobalt, tin, etc. Thus, the number of alloys that are available for the various processes is not very limited in modern AM technology as it was at in earlier stages. Sometimes postprocessing, secondary operations, and finishing are required to remove (or dissolve) supports that are used to support overhanging features during construction. Improvement is required to eliminate the necessity of finishing process. Thus, there is still a lot of work and research to be accomplished before AM technologies become standard in the manufacturing industry because not every commonly used manufacturing material can be handled. AM processes take the information from a CAD file that is later converted to a stereolithography (STL) file. The increasing development and the successful results of AM since its early use predict that this technology has a significant place in future and current manufacturing techniques. In practice, the terms 3D printing and AM may be used interchangeably.

CHAPTER TWO

Fabrication In this chapter, a short description of the processes mentioned in Chapter 1 is presented with some basic equations applicable to these processes.

2.1 Fused deposition method (FDM) The FDM method is applied widely for producing plastic parts. It is one of the several additive manufacturing processes and was about the first rapid printing technology invented about two to three decades ago. The technology became known in the 1980s and was used to print plastic objects. Due to the flexibility of this technology being capable of producing thermoplastic objects of almost any shape and geometry, the aerospace industry is using FDM for a variety of applications. A further advantage of this revolutionary technology is that the objects can be manufactured in one piece without the need to drill, weld, or attach individual components together. In the conventional manufacturing technique, producing an intricate object in one piece is not feasible. A schematic outline of an FDM production cycle is illustrated in Fig. 2.1. The FDM which is a typical rapid printing (RP) process produces threedimensional objects which are formed from a computer-generated design. In Fig. 2.2, a schematic illustration of the FDM extrusion and deposition process is shown. The temperature-controlled extruder forces out a thermoplastic filament and deposits the semi-molten material (polymer) onto a substrate (platform) layer by layer. The filament is moved between two rollers as shown in Fig. 2.2 and the semi-molten polymer is thrust out to the platform. On finishing each layer, the base platform is lowered for the start of the next layer and this is repeated several times until fabrication of the designed project is completed as a three-dimensional object. The deposition path of every deposited layer depend on the material used, the fabrication conditions, and the application of the designed part. A schematic representation of the FDM process is illustrated schematically in Fig. 2.3. In summary, FDM has three fabricating steps: (1) preprocessing, (2) processing, and (3) post-processing. In (1) a CAD model is constructed in stereolithography (STL) file format for the FDM process, (2) in the FDM Additive and Traditionally Manufactured Components https://doi.org/10.1016/B978-0-12-821918-8.00002-4

© 2020 Elsevier Inc. All rights reserved.

3

FIRST LAYER DEPOSITION I

z

II

heating chamber

LAST LAYER DEPOSITION III

3D PRINTED OBJECT IV

filament

y x

LOWERING OF THE BUILDING PLATFORM

tip

z y

post-process

x

3D PRINTED FINISHED OBJECT

Fig. 2.1 Outline of a FDM production cycle. (Melocchi, A., Parietti, F., Loreti, G., Maroni, A., Gazzaniga, A., Zema, L., 2015. J. Drug Delivery Sci. Technol. 30, 360.)

5

Fabrication

filament

X-direction

FDM head

Y-direction

platform Z-direction

FDM head PCL filament rollers

liquefier

temperature control

x-y axes nozzle tip PCL extrudate

scaffold

platform z-axis

Fig. 2.2 A schematic diagram of the FDM extrusion and deposition process. The nozzle tip could be changed into different sizes for different RW and thickness. The layers of roads were fused together upon solidification to form a 3D structure. RW stands for road width. (Zein, I., Hutmacher, D.W., Tan, K.C., Teoh, S.H., 2002. Biomaterials 23, 1169. With kind permission of Elsevier.)

6

Additive and traditionally manufactured components

ComputerAided Design (CAD) model

export

.stl (StereoLithography) format

import

Stratasys’ QuickSLice (QS) software

slicing & setting parameters

Physical model FDM process FDM machine

download data

.sml (Stratasys Machine Language) format

create tool paths or

.slc (SLiCe) format

Fig. 2.3 Summary of a basic FDM process. Step 1: Import of CAD data in stl (StereoLithography) format into QuickSlicet. Step 2: Slicing of the CAD model into horizontal layers and conversion into an slc (SLiCe) format. Step 3: Creation of a deposition path for each layer and conversion into an sml (Stratasys Machine Language) format for downloading to the FDM machine. Step 4: FDM fabrication process using a filament modeling material to build an actual physical part in an additive manner layer by layer. (Zein, I., Hutmacher, D.W., Tan, K.C., Teoh, S.H., 2002. Biomaterials 23, 1169. With kind permission of Elsevier.)

machine process, the layers are formed until completion of the model, and (3) in this stage, the model and any supports are removed by washing or stripping away. Finally, the surface of the model is cleaned by finishing. Thus, the key elements in a typical process include the feed mechanism of the material, liquefier, and print nozzle. Relationships of a pinch roller feed mechanism (Turner et al., 2014; Pandey and Pradhan, 2018) according to Bellini et al., 2004 are presented below. A pinch roller is seen in the lower part of Fig. 2.2. The linear velocity of the filament v can be expressed by v¼

Q WH

(2.1)

where Q is the constant volumetric flow rate, W is the slice thickness, and H is the slice thickness. Assuming a perfect adhesion between filament and roller, the feed velocity can be expressed as v ¼ wr Rr

(2.2)

where wr is the angular velocity and Rr is the radius of the rollers, respectively. The force, F to push the melt through the liquefier is determined from the pressure drop (ΔP).

7

Fabrication

F ¼ ΔPA

(2.3)

A being the cross-sectional area of the filament assumed to be equal to the cross-sectional area of the liquefier. The required torque (Γ ) is given as F Γ ¼ Rr 2

(2.4)

The power, Pmo is given by the product of the angular velocity and the torque as Pmot ¼ wr Γ

(2.5)

In Eq. (2.5) it is assumed that two motors provide the power to the roller mechanism. When the compression reaches a critical power, filament can buckle which is the most common failure. Euler’s buckling analysis can approximate the critical pressure, Pcr Pcr ¼

π 2 Edf2 16Lf2

(2.6)

where E is the elastic modulus, df is the filament diameter, and Lf is the filament length measured from the rollers to the entrance of the liquefier. The size of the printed material depends on the amount of melt in the liquefier. The temperature, the viscosity, and surface energy of the melt determine the feed rate which controls the flow through the liquefier. Temperature, viscosity, and surface energy of the melt provide the effective parameters to join and bond neighboring printed polymer units (parts). It is important to consider the mentioned parameters, namely, temperature, viscosity, and surface energy of the melt on the bonding of the printed polymer units.

2.1.1 Melt properties The viscous behavior of the melt (polymer in the example) is expressed by a power law (for extrusion, it is a shear-thinning process; recall the shearthinning process allows flow under proper shear stress but recovers back to the solid after removal of the stress) as η ¼ K ðγ_ Þn1

(2.7)

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Additive and traditionally manufactured components

η is the viscosity, γ_ is the shear rate, K and n are power-law fit parameters. The change in viscosity with temperature is the product of a temperature and a shear-rate–dependent terms at some reference temperature given as η ¼ H ðT ÞηT0 ðγ_ Þ

(2.8)

An Arrhenius relation is used for analyzing the flow behavior of a polymer melt and the viscosity variation with temperature as 
 1 1 (2.9)  H ðT Þ ¼ exp α T T0 It has been assumed that the heat capacity of melt, Cp is constant, however, it is not, and is known to change significantly at the glass transition temperature (Tg) for amorphous polymers. The specific heat of melt varies with respect to temperature by the following relations (Bellini, 2002): T < 300K Cp ¼ 4:4T + 58 T > 300K Cp ¼ 105T + 16:68

(2.10)

2.1.2 Liquefier The liquefier has a major effect on the melting behavior. Most current liquefiers are cylindrical and contain three regions as seen in Fig. 2.4

material flow L1

I

L2

II

D1

b

III D2

Fig. 2.4 Liquefier divided into three zones for modeling. (Turner, B.N., Strong, R., Gold, S.A., 2014. Rapid Prototyp. J. 20, 192. With kind permission of Elsevier.)

9

Fabrication

The analytical models assume either a constant heat flux or a constant wall temperature. According to Bellini, in the constant heat flux model, the required heat flux to the liquefier can be approximated by 0 1 B ρvA C C q ¼ mc _ p ðT  Ti Þ ¼ B @
D Acp ðT  Ti Þ 2π L 2

(2.11)

m_ is the mass flow rate of the polymer through the liquefier, cp is the heat capacity of the polymer, D is the diameter, L is the length of the liquefier, and T and Ti are the temperatures of the polymer at exit and entrance of the liquefier, respectively.

2.1.3 Heat convection Temperature gradients within a liquefier were evaluated by Yardimci et al. (1997) using finite element analysis. Heat convection in the liquefier entrance and nozzle exit was examined. Convective heat transfer coefficients, h of 10 W/m2 K at the entrance and 100 W/m2 K at the nozzle exit were assumed by Yardimci et al. (1997) For the location of the melt front, a two-dimensional axisymmetric steady-state advection-conduction model was assumed for the filament heating and a constant wall temperature and plug flow (Yardimci et al., 1997). The nondimensional solution to this problem is given as θ¼2

∞ X n¼1

  F 0 ðλn r 0 Þ exp λ2n z0 F 0 ðλn Þ ¼ 0 λn F 1 ðλn Þ0

(2.12)

where θ, r0 , and z0 are the dimensionless temperature, radius, and length, respectively, defined as θ¼

T  T0 0 r αz r ¼ 0 , z0 ¼ 2 T2  T 0 rF θrF

(2.13)

where T is the filament temperature at some point (r, z), T0 is the wall temperature at the entrance of the liquefier (z ¼ 0), rf is the filament radius, α is the thermal diffusivity in the filament, λn are the roots of zero-order Bessel function of the first kind F 0 , and F 1 is the first-order Bessel function of the first kind. Constant thermal properties of the filament are assumed in the solution and any gap between the filament and liquefier wall is neglected. The location of the melt front can be calculated when the temperature of

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Additive and traditionally manufactured components

the center of the filament is equal to the melt temperature. This is the condition of solving for z0 that occurs when r ¼ r0 and T ¼ Tm (Yardimci et al., 1997). For the analysis, the liquefier is divided into three sections, as shown in Fig. 2.4.

2.1.4 Pressure drop estimation For calculating the optimum feed rate at the liquefier inlet, the pressure drop inside the liquefier has to be known. Prediction of the melt velocity profile, the pressure drops and the shear stress profile can be performed from a momentum balance in the liquefier. With the help of energy balance, the equation temperature profile of polymer melt inside the liquefier can be determined. Based on the work of Bellini and others (Ramanath et al., 2008), solutions to the momentum balance equations were developed for extrusion dies (Michaeli, 2003) for cylindrical, conical, and cylindrical shapes corresponding to regions I, II, and III (Fig. 2.4). The assumptions considered during the dynamic analysis are the melt is incompressible, no-slip boundary conditions apply at the walls of the liquefier that the flow is fully developed, steady-state and laminar conditions apply. The pressure drops in each section of the liquefier according to this model are given, respectively, by 0 11 = m 
 
1 = v mB m + 3 C 1 1 ΔP1 ¼ 2 L1 exp α  @ m + 1 A ϕ T Tα D1 = 2

(2.14)

! !1 = !
2 m 2m 1 1 D1 m+3 β   ΔP2 ¼  ð m + 3 Þ2 3 3 = = 2 3 tan =2 D2 m D1 m 
 1 1 (2.15) exp α  T Tα 0


2 11 =m D 

1 = Bðm + 3Þ 1 C Δ mB 1 1 2 C ΔP3 ¼ 2L3  B
m + 3 C exp α @ A ϕ T Tα D2

(2.16)

2 The dimensions of L1, L3, D1, and D2 are illustrated in Fig. 2.4, β is the nozzle angle of the liquefier (indicated in Fig. 2.4), and m and ϕ are power-law

11

Fabrication

fit parameters. Thus, the total pressure drops, and ΔP inside the liquefier is the sum of the pressure drops in different sections, ΔP ¼ ΔP1 + ΔP2 + ΔP3

(2.17)

Further analysis for obtaining the finished product is associated with road deposition, spreading, and bonding.

2.1.5 Layer deposition and stability The volumetric flow of the polymer melt extruded from the nozzle exit (Fig. 2.4) lands on a build sheet or on a previously deposited layer of polymer. The analysis is according to the Hagen-Poiseuille flow. The volumetric flow rate, Q, is given by


 D2 4 π 2 Q¼ 8ηL2

(2.18)

The cross-sectional area, A of the bead is inversely related to the velocity of the print head vprint A¼

Q vprint

(2.19)

At sufficiently large head velocity the layer is unstable and discontinuous. The deposition at maximum velocity for which the layer cross section remains stable and continuous is estimated (Middleman, 1995) by axisymmetric flow (a flow pattern is said to be axisymmetric when it is identical in every plane that passes through a certain straight line. The straight line in question is referred to as the symmetry axis) νprint
10

1
deqm no continuous track of liquid is formed. Eq. (2.46) is valid only when w > deqm, but if w < deqm the liquid track has nonparallel sides as seen

d

θ w

p

Fig. 2.26 Schematic illustration of the coalescence of individual drops to form a track or liquid bead with a uniform cross section of a circular sector. (Derby, B., 2011. J. Eur. Cer. Soc. 31, 2543. With kind permission of Elsevier.)

31

Fabrication

in Fig. 2.27b. The maximum spacing of drops, pmax to produce parallel-sided liquid bead can be obtained by setting w ¼ pmax pmax ¼

3β

2



θeqm

= sin 2 θ

2πd0    eqm  cosθ eqm = sinθ eqm

(2.47)

Fig. 2.27 shows the behavior of inkjet-printed tracks as the drop spacing reduces under the conditions indicated in the figure. The conditions for a stable line are given by UT∗ > gðp∗ , θadv Þ

(2.48)

with UT∗ ¼

UT η γ

(2.49)

Fig. 2.27 Four morphologies possible when individual drops are printed onto a surface at regularly spaced intervals: (a) drops are spaced p > deqm: no interaction occurs, (b) pmax < p < deqm: a continuous track is formed but contact line pinning results in an irregular edge, (c) p < pmax: parallel-sided track is formed, (d) when drop spacing is below a threshold determined by both contact angle and printing speed, a bulge instability develops. (Derby, B., 2011. J. Eur. Cer. Soc. 31, 2543. With kind permission of Elsevier.)

32

Additive and traditionally manufactured components

U∗T is a dimensionless traverse velocity and g(p*,θ) is a dimensionless drop spacing. The function g(p*,θ) is related to the inverse drop spacing and the contact angle as shown in Fig. 2.28. A graphical representation of the range of conditions under which stable parallel lines can be printed from a train of discrete drops is shown in Fig. 2.29. Droplet drying by evaporation process can be estimated from the relation Ma ¼

Δγr ηD

(2.50)

where Ma is the Marangoni number and a dimensionless parameter, Δγ is the difference in surface tension between the two pure solvents, r is a characteristic length assumed to be the radius of the spread drop on the substrate, η is the liquid viscosity, and D is the solute diffusion coefficient. The quality of the printed object depends on the interaction of the droplets and the drop drying process. The use of a proper evaporating solvent in printing ceramics is important to ensure maximum solid density prior to sintering. Zirconia - ref [35] Ag Nanoparticles - ref [32] PEDOT/PSS 1 - ref [34] PEDOT/PSS 2 - ref [34]

100 10-1

STABLE LINES

U*T

10-2 10-3 10-4

BULGE INSTABILITY

10-5 10-5

10-4

10-3 10-2 g(p*, qadv)

10-1

100

Fig. 2.28 The bulge instability is observed below a critical drop spacing that depends on the traverse speed of the printer. The criterion for the onset of the instability in terms of the parameter g(p*, θ) and a dimensionless velocity is shown as the solid line, with experimental data from a range of sources superimposed. Black symbols indicate well-formed lines with parallel sides, and white symbols indicate the conditions under which unstable bulges appear on printed lines. (Derby, B., 2011. J. Eur. Cer. Soc. 31, 2543. With kind permission of Elsevier.)

33

Fabrication

15°

10-1

Irregular Line Limit (Changes with Contact Angle)

30°

100

U*T

10-2 10-3 10-4

10-5 10-5

60° 90° 45° 75° Transverse Velocity Limit Region of Stability it

y ilit

Lim

tab

ns

eI ulg

B

10-4

10-3 10-2 g(p*, qadv)

10-1

100

Fig. 2.29 Graphical representation of the range of conditions (drop spacing and printer speed) under which stable parallel lines can be printed from a train of discrete drops. The diagonal line represents the onset of the bulging instability and a minimum drop spacing. The vertical dashed line represents the maximum drop spacing to form a parallel-sided track. The horizontal dotted line indicates an upper limit for printer speed that represents the mechanical limitations of a given printer system and the value is shown in this figure is purely arbitrary. (Derby, B., 2011. J. Eur. Cer. Soc. 31, 2543. With kind permission of Elsevier.)

2.4 Stereolithography (SLA) Inventors in several countries (Japan, France, and USA) filed patents on three-dimensional printing of polymers but the inventor Chuck Hall (USA) coined the term stereolithography in the 1980th. Stereolithography is also an additive manufacturing process that uses ultraviolet light to cure photosensitive polymers. The ultraviolet source can be a laser focusing its light onto a vat of photopolymer resin. As in the earlier mentioned AM fabrication, the process is performed with the help of CAM/CAD software. The ultraviolet light solidifies the photopolymer sensitive plastic (resin) forming a single layer of the 3D object. The printing by the stereolithographic technique—also known as an optical fabrication—is among the earliest 3D manufacturing methods and is widely used for various materials such as: (i) Standard resin (ii) Engineering resin (ideal for high-temperature applications due to toughness and flexibility) (iii) Castable resin

34

Additive and traditionally manufactured components

(iv) Dental resin (v) But also, for the production of ceramic components. The ceramic component is incorporated in a proper solvent. The technique is also used for creating models, prototypes (still in early design), patterns, and production of parts in a layer-by-layer fashion. The process induced by light causes photopolymerization and linking of chains of molecules forming the polymers which make up the three-dimensional solid. The technique is a very important method to produce biocompatible materials for medical components and is successfully used for medical purposes such as orthopedic applications (finger joints and parts of bones) and tissue engineering (artificial skin in treating burns). The products are of high accuracy and smooth so that in many cases no surface finish is required. Besides the highest accuracy of the technique and the smoothest surface finish of all 3D processes, it is capable of producing very detailed small parts and very large parts with great resolution and virtually without shrinkage or warping. Fundamental relationships for stereolithography were developed by Jacobs (1992) applying the following assumptions: (i) The polymer obeys the Beer-Lambert law of exponential absorption (ii) The laser radiation distribution is Gaussian (iii) The transition from the liquid to the solid-state occurs at the gel point From Beer-Lambert law
 z H ðx, y, zÞ ¼ H ðx, y, 0Þ exp  (2.51) Dp In defining the coordinate system the laser is scanned along the x-axis, y is orthogonal to the laser scan axis and the z-axis is normal of the x-y plane (the resin-free surface where z ¼ 0) and extends downward into the resin. Thus some arbitrary point Q (x, y, z) defines the origin. The projection of Q onto the resin surface is Q0 (0, y, 0), showing that the origin is arbitrarily selected such that both x and z are zero. In Eq. (2.51), H(x,y,z) is the irradiance at any arbitrary point, H(x,y,O) is the surface irradiance at any point x,y,0, and Dp is the penetration depth. For a Gaussian distribution as assumed above in (ii) Dp is defined as the depth of resin which reduces the irradiance to 1/e of the surface irradiance (37%). Then one can write instead of Eq. (2.51),   2r 2 H ðx, y, 0Þ ¼ H ðr, 0Þ ¼ H0 exp (2.52) W02

35

Fabrication

W0 is the radius of the Gaussian beam defined at the 1/e point (13.5% of the peak irradiance, H0). It is convenient to use cylindrical coordinates because the Gaussian function is circularly symmetric. From the laser power incident P1 on the polymer surface, H0 can be determined as the integral of the laser irradiance distribution over the entire surface, namely from r ¼ 0 to r ¼ ∞. Z

∞

P¼

Z

∞

H ðr, 0Þ2πrdr ¼ 2πH0

0

0



 2r 2 exp  2 rdr W0

(2.53)

The expression in the exponential is defined as
2 2r 2 W0 du u  2 or rdr ¼ W0 4

(2.54)

Substituting (2.54) in (2.53) one obtains π PL ¼ W02 H0 2

Z 0

∞

π exp ½udu ¼ W02 H0 2

(2.55)

From (2.55) H0 ¼

2P1 πW02

(2.56)

Coming back to Eq. (2.53) and substituting (2.56) and then into Eq. (2.51) we obtain the Gaussian laser irradiance (cylindrical coordinates),


   2PL z 2r 2 H ðr, zÞ ¼ exp   2 πW02 Dp W0

(2.57)

Now going to the exposure function E referring to the actinic exposure. Note that the actinic flux describes the number of photons incident at a point, while the irradiance describes the radiant energy crossing a surface. Z E ¼ Hdt (2.58) The velocity VS of the laser along the x-axis is VS ¼ dx=dt or dt ¼ dx=VS

(2.59)

36

Additive and traditionally manufactured components

Substitute relation (2.59) and Eq. (2.57) into (2.58) to get
  Z ∞   2PL z 2r 2 E ðr, zÞ ¼ exp  exp  2 dx πW02 VS W0 Dp ∞

(2.60)

The exponential term is outside the integral since z is not a function of x and thus considered as a constant. Also since (Pythagoras), r 2 ¼ x2 + y2

(2.61)

      2r 2 2x2 2y2 exp  2 ¼ exp  2 + exp  2 W0 W0 W0

(2.62)

one can write

Substituting Eq. (2.62) into Eq. (2.60) and again bringing outside the integral the exponential factor (as it is not a function of x) and since the integral is symmetric about x (and therefore the total integral is twice the value from 0 to infinity) one can write
Eðy, zÞ ¼

  Z ∞   4PL z 2y2 2x2 exp  dx  exp  πW02 VS W02 Dp W02 0

(2.63)

define v as pffiffiffi 2x v W0

(2.64)

differentiate it to obtain
pffiffiffi
 W0 2 dx or dx ¼ pffiffiffi dv dv ¼ W0 2

(2.65)

One can evaluate the integral in Eq. (2.60) by substituting from Eqs. (2.64) and (2.65)  
Z ∞ Z ∞

 2x2 W0 (2.66) exp  2 dx ¼ pffiffiffi exp v2 dv W0 2 0 0 The integral is related to the error function. It is found that Z 0

∞

 exp v2 dv ¼

Z pffiffiffi π 2

(2.67)

37

Fabrication

Now substituting Eqs. (2.66) and (2.67) into Eq. (2.63) we obtain  
1:2  2 PL z 2y2 exp  + 2 (2.68) E ðy, zÞ ¼ W0 vS π Dp W0 The exposure is at its maximum value E ¼ Emax when y ¼ 0 (the laser scan axis) and z ¼ 0 (namely on the polymer or the resins surface). Emax is thus given by
1=2 2 PL (2.69) Emax ¼ W0 vs π Eq. (2.69) shows that the maximum actinic laser exposure is proportional to the laser power. Inversely proportional to the product of the beam radius and the scan velocity and the proportionality constant is (2/π)1/2 (¼0.7979).

2.4.1 The state of the resin (photopolymer) When E < Ec. Ec being a critical value, the resin is liquid and when E > Ec the resin partially polymerizes, but when E ¼ Ec the resin is at the gel point (the gel point refers to an abrupt change in the viscosity of a solution containing polymerizable components) corresponding to the liquid undergoing transition to a solid phase. If at the gel point y and z of Eq. (2.68) are y ¼ y∗ and z ¼ z∗, all points inside this locus are partially solidified, while outside this boundary is still liquid. The cross-sectional shape of a single laser cured photopolymer after inserting the values of Ec, y∗, and z∗ into Eq. (2.68) one obtains  2 
1=2  2y∗ z∗ 2 PL ¼ (2.70) + exp W02 Dp W0 vS Ec π Taking natural logarithm of (2.70) and substituting Emax from (2.69) one can write   2y∗ 2 z∗ Emax (2.71) + ¼ ln W02 Dp Ec Rewriting Eq. (2.71) in the form of Ay∗ + Bz∗ ¼ C where

  2 1 Emax A¼ 2 B¼ C ¼ ln Ec W0 Dp

(2.72)

38

Additive and traditionally manufactured components

This is the equation of a parabolic cylinder in three dimensions whose axis is the x-axis which is by the above definition the laser scan axis.

2.4.2 The maximum cure depth Define the cure depth by Cd. From Eq. (2.71) z∗ ¼ zmax ¼ Cd when y* ¼ 0. Setting these values in Eq. (2.71) we obtain   Emax Cd ¼ Dp ln (2.73) Ec The meaning of this relation is (i) The cure depth is the natural logarithm of the maximum actinic laser exposure. (ii) The relation of (2.73) is a straight line on a semilogarithmic plot of Cd vs. Emax. (iii) (it is known as the working curve). (iv) The slope of the working curve is the penetration depth, Dp of the resin. (v) Since ln(1) ¼ 0 the intercept of the curve is equal to the critical exposure, Ec of the resin (at the laser wavelength). (vi) The slope and the intercept are independent of the laser power, spot size, W0 or the laser scan velocity vS (since Dp and Ec are just numbers).

2.4.3 The cured line width From Eq. (2.71) the maximum cured linewidth, LW will occur at the resin surface where the parabolic figure has its greatest width. It is obtained by setting y∗ ¼ yð max Þ ¼ LW =2 where x∗ ¼ 0. Thus pffiffiffi LW ¼ 2W0





Emax ln EC

 1=2 (2.74)

Substituting for ln[Emax/Ec] from Eq. (2.73), Eq. (2.75) is obtained for the cured linewidth, sffiffiffiffiffiffiffiffi 2Cd (2.75) LW ¼ W0 Dp

39

Fabrication

Thus, the relation shows: (i) The cured linewidth is directly proportional to the laser spot size at the plane of the resin surface. (W0 is the radius of the spot, not diameter). (ii) The linewidth is also proportional to the term in the square root (ratio of the cure depth to the resin penetration depth). (iii) The linewidth also depends on the resin penetration depth.

2.4.4 Laser scan velocity Eq. (2.73) can be rewritten as 

Cd Emax ¼ EC exp Dp

 (2.76)

Substituting for Emax from Eq. (2.69) a relation for Emax is obtained as  
1=2  2 PL Cd ¼ EC exp Emax ¼ W0 VS Dp π

(2.77)

Solving for the laser velocity VS we obtain  
1=2  2 PL Cd exp  VS ¼ W0 EC Dp π

(2.78)

The meaning of the relation is (i) The laser velocity is directly proportional to the laser power, the higher it is the faster the scan speed for a given resin for a specific laser spot size and cure depth. (ii) The laser scan velocity is inversely proportional to its spot size (for a given resin for a specific laser spot size and cure depth). (iii) The scan velocity decreases exponentially with the increase of the ratio in the exponential term. This explains why increased cure depth for a certain polymer draws more slowly than shallow cure depth. (iv) The constant of proportionality is a pure number (2/π) ¼ 0.7979. The above relations have been experimentally confirmed in stereolithographic AM. The relations presented are the foundation of this technology and provide basic physical understanding of the interaction of photopolymers with actinic photons. The theoretical predictions of the model described are in good agreement with much of the experimental results.

40

Additive and traditionally manufactured components

2.5 Direct energy deposition (DED) The method is defined as an additive manufacturing processing in which the used thermal energy melts (or fuses) materials as they are deposited. The energy source for the process can be a laser, electron beam, or plasma arc, which is focused on the material to melt it. A wide range of materials such as polymers, ceramics, and a variety of metals and alloys including ferrous alloys can use DED additive manufacturing but DEDs special use is repairing and maintaining structural parts. The machine typically consists of a nozzle, which can move in many directions and is not fixed to a specific axis, therefore the material can be deposited from any angle and melted. The process is quite similar to the extrusion process. The machine consists of several arms with the nozzles that can move around an object. As in AM, the material is deposited from the nozzle onto the existing surfaces of the object. In case of metal, it can be supplied as powder or wire. During the process, material is added layer by layer which solidifies, creating new objects or repairing existing objects. A thermomechanical model for the DED process is presented further, which was implemented for Ti-6Al-4V. Starting with a heat transfer energy balance.

2.5.1 Thermal model The heat energy balance (following Heigel et al., 2015) is ρCp

dT ¼ r  qðr, tÞ + Qðr, tÞ dt

(2.79)

where ρ is the material density, Cp is the specific heat capacity, T is the temperature, t is the time, Q is the heat source, r is the relative reference coordinate, and q is the heat flux vector, calculated as q  krT

(2.80)

Here k is the thermal conductivity of the material. Above 850°C, the thermal properties are assumed to be constant for Ti-6Al-4V and the density of 4.43  103 kg/m3 is assumed to be T independent over a temperature range from 1600°C to 1670°C. The laser heat source is described by the double ellipsoid model, accordingly the heat source, Q can be expressed as

41

Fabrication

" # pffiffiffi 6 3Pηf 3x2 3y2 3ðz + vW tÞ2 pffiffiffi exp  2 + 2 + Q¼ a b c2 abcπ π

(2.81)

P is the laser power, η is the laser absorption coefficient, x, y, and z are local coordinates, while f, a, b, and c define the volume over which the heat source is distributed. The value η is found to be 45%. The volume is defined such that the circular heat source has a radius equal to half of the deposition track width and applied to a depth of 0.9 mm, which results in a melt pool depth to radius ratio of 0.6. Heat loss occurs on all the surfaces by convection, qconv and radiation, qrad. Radiation is defined by the Stefan-Boltzmann law:   4 (2.82) qrad ¼ εσ Ts4  T∞ where ε and σ are the surface emissivity and the Stefan-Boltzmann constant. Ts and T∞ represent surface and ambient temperatures, respectively, and the emissivity is temperature independent (¼0.57). The surface heat loss is defined as   4 qconv ¼ h Ts4  T∞ (2.83) h is the coefficient of convection. A detailed evaluation of the above model for better understanding of experimental results (Kumar et al., n.d.) is expanded and presented below: The Gaussian profile is given as
 2r 2 P ðr Þ ¼ P0 exp  2 (2.84) r0 P is the calculated power at radius r, P0 the given laser power, r0 the radius of the laser beam, and r the current radius. The heat conduction rate is proportional to the heat transfer area and the temperature gradient and is given as qcond 52kA

∂T ∂x

(2.85)

As indicated earlier, k is the thermal conductivity (W/m K). Assuming planar layers and the three-dimensional heat conduction (Cartesian coordinates) is given as

42

Additive and traditionally manufactured components

ρCp


2  ∂T ∂ T ∂2 T ∂2 T + + ¼k ∂t ∂x2 ∂y2 ∂z2

(2.86)

Here ρ is the density. As mentioned above convection occurs to the surrounding (assumed to be air) at a rate where Ta is the air temperature. βρconv ¼ hAðTs  Ta Þ

(2.87)

h in units (W/m2 K) and A is the face of a solid area. The radiation losses can be obtained from Stefan-Boltzmann law providing the total emissive power of a blackbody, Eb as Eb ¼ σΤ 4

(2.88)

T is the blackbody temperature and σ the Stefan-Boltzmann constant is 5.67  108 W/(m2K4). When a body surface area, A, of temperature Ts, is immersed in a medium with ambient temperature Ta, the heat rate radiated by the body is given by [see Eq. (2.83)].   (2.89) qrad ¼ σA Ts 4  Ta 4 The surrounding medium is air at temperature Ta ¼ 298 K. The heat absorbed or released during the phase change of an element, Lf with mass m is qlatent

heat

¼ mL f

(2.90)

Now let us consider the energy balance by the following boundary condition: k

 4  ∂T ηPlaser 4  h ð T  T Þ  εσ T  T ¼ s a s a πR2 ∂n

(2.91)

in Eq. (2.91) laser irradiation, convective losses, and radiative losses are taken into consideration; η is the laser absorption coefficient, Ppower is the power of the laser and is obtained from Eq. (2.84), R is the radius of the laser spot, n is the thermal vector at the local interface, and ε is the emissivity. The losses of radiation on the side and bottom surfaces are assumed to be negligible. The boundary condition at the sides and bottom are K

∂T + hðTs  Ta Þ ¼ 0 ∂n

(2.92)

A fluid model and additional details including experimental parameters are also presented by Kumar et al. (n.d.) and more on the subject can be obtained

43

Fabrication

by consulting his work. The work deals specifically with the effect of laser exposure on Ti-6Al-4V alloy. A comparison of simulated and experimental results shows that the model presented by the above equations predicts the temperature profile and solidified metal profile, respectively.

2.6 Laminated object manufacturing (LOM) Laminated object manufacturing is an additive manufacturing technology that can be effectively used in addition to a paper for a range of materials such as ceramics, metals polymers, and composites with limitations dictated by the material and object to be formed. A typical LOM contains components such as a portable mirror, heater-roller—and in the case of paper objects—wastepaper take-up roller, paper feed roll, and laser. A typical LOM process is illustrated in Fig. 2.30. The first LOM system was developed

Laser Optics

X-Y positioning device Laminating roller

Layer outline and crosshatch

Sheet material Part block

Platform

Take-up roll

Material supply roll

Fig. 2.30 LOM process. (Liao, Y.S., Chiu, Y.Y., 2001. Int. J. Prod. Res. 39, 3479. With kind permission of Taylor and Francis.)

44

Additive and traditionally manufactured components

by Helisys of Torrance, CA, USA. As indicated the main components in a LOM system are a feed mechanism (that advances a sheet over a build platform), a heated roller to apply pressure to bond the sheet to the layer below, and a laser to cut the outline of the part in each sheet layer. An adhesivecoated sheet is on the top of the previous sheet for bonding (under pressure and heat). The purpose of the laser is to cut the outline (of the paper) of the part into each layer. After completion of each cut, the platform (see Fig. 2.30) lowers to a certain depth, usually equal to the sheet thickness and another sheet is advanced on top of the earlier deposited layers. The heat and pressure are applied after the platform rises to bond the new layer to the previous ones. After laser cutting the outline, the process is repeated until the entire component or part is completed. Thus, this class of AM is also an automated fabrication process to produce a three-dimensional object built directly from a surface or solid CAD file by sectioning electronically a designed object. The building of the object is from the bottom up one layer at a time. Note that after a layer is cut, the extra material remains in place to support the part during the build. The process is applicable not only for paper but also for plastic and metal laminate coated with an adhesive such as polyethylene. LOM is especially suited to larger parts, parts with larger wall thickness, and pattern making. Other applications are for using in silicone rubber molding, vacuum forming, and fiberglass prototyping. Bonding of layers is an essential part of the LOM process. In Fig. 2.31, stages of the LOM process are indicated schematically. A possible chemical reaction in intermediate layer is suggested in (C) as indicated in Fig. 2.31. The process build time in LOM manufacturing is important and an algorithm for estimating the process by Kechagias et al. (2004) is presented

A

A G

B

a)

G

B

b)

external heating

A B

reagent

c)

Fig. 2.31 Schematic representation of some stages of LOM process: (a) process without intermediate layers formation; (b) process with intermediate layers formation; and (c) post-processing with external heating. (Knyazeva, A., Travitzky, N., 2018. 3rd Inter. Conf. on Rheology and Modeling of Mater. IOP Publishing, IOP Conf. Ser. J. Phys. Conf. Ser. 1045, 012020; Open access.)

45

Fabrication

below: the total process build time, Ttotal is the sum of the time for the layer preparation, TTPL and the time for the layer processing, TLP given as Ttotal¼ TTPL + TLP   TTPL ¼ n TPlatform Down + TFeeder + TPlatformup + THeater + TDelay TLP ¼ TLaseron + TLaseroff + TLaseridle

(2.93) (2.94) (2.95)

The number of layers is given by n¼

Z LT

(2.96)

where LT is the layer thickness and Z is the coordinate difference Zmax  Zmin.. The times indicated in Eq. (2.94) are given as follows: TPlatformDown ¼

DPdown PS

(2.97)

DP–down is the platform downward distance (see Fig. 2.30) and PS is the platform speed. TFeeder ¼

X + 2DX + AM PS

(2.98)

Here X is the coordinate difference Xmax  2Xmin, Dx is the gap between the inner and outer boxes that surrounds the part layer, AM is the advance margin, i.e., the distance between two successive outer boxes and PS is the feeder speed. TPlatgormup ¼

DPup PS

(2.99)

DP–up and PS are the platform’s upward distance and PS is its upward speed. THeater ¼

2 HM ðHRM + X + 2DX + HLM Þ + HS HS

(2.100)

where HRM is the heater right margin, DX is the gap between the inner and the outer boxes that surround the part layer, HLM is the heater left margin, HM is the heater margin, and HS is the heater speed. TDelay which is the time of software responsible for the platform upward adjustments is not calculated by the slicing software.

46

Additive and traditionally manufactured components

TLaser-on of Eq. (2.95) is given as   ðX + DX + Y + DY Þ TLaser0n ¼ 4n LS BoxPerimeter   SSTL  SPlat + LT  LS PartPerimeter 20 3 1 VSTL
 ðn  1Þ  XY  1 1 7 6B LT C + 4@ + 5 A LS HX H Y 






Hatch

ði  XY  SFC Þ 1 1 + (2.101) LS HXPC YPC FineCrossHatch 
   n  ðX + Y Þ VSTL 1 1 ¼ + + ðLT  LSÞ HX HY Part LS Hatch   ð2  OffsetÞ + n (2.102) LS Offset +

TLaseroff

and


 1 X Y n + TLaseridle ¼ 1 + 4 HY HX

(2.103)

In the above relations, Dx and DY are the gaps between the inner and outer boxes that surround the part layer in both the X and Y directions, HX and HY, HXFC and HYFC are the hatchings and fine cross-hatching spacings in the X directions and the Y directions, respectively, i is the number of crosshatched layers, LS is the layer speed and LT is the layer thickness, n is the number of layers involved, SSTL, SFLAT, and SFC are the part surface, the flat surface, and the total fine cross-hatched surface, respectively, VSTL the part volume and X (Xmax  Xmin) and Y (Ymax  Ymin) is the coordinate differences. Eqs. (2.93)–(2.103) represent the basic build-time algorithm of the LOM process. Recall that Fig. 2.30 is a schematic illustration of the LOM process where layers are gradually added on a platform (in this case each layer corresponds to the specially coated sheet of paper), while a heater glues each layer on the previous layer. The laser beam is used to cut the part of cross section on each layer to be glued together to build up the object of interest.

Fabrication

47

References Bellini, A., 2002. Fused Deposition of Ceramics: A Comprehensive Experimental, Analytical and Computational Study of Material Behavior, Fabrication Process and Equipment Design (PhD Dissertation). Drexel University, Philadelphia, PA. Bellini, A., G€ uc¸eri, S., Bertoldi, M., 2004. J. Manuf. Sci. Eng., Trans. ASME 126, 237. Bidare, P., Maier, R.R.J., Beck, R.J., Sheph, J.D., Moore, A.J., 2017. Addit. Manuf. 16, 177. Crockett, R.S., 1997. The Liquid-to-Solid Transition in Stereodeposition Techniques (PhD Thesis). University of Arizona. Derby, B., 2011. J. Eur. Ceram. Soc. 31, 2543. Heigel, J.C., Michaleris, P., Reutzel, E.W., 2015. Addit. Manuf. 5, 9. Jacobs, P.F., 1992. Fundamentals of Stereolithography. 3D Systems, Inc, Valencia, CA. Kechagias, J., Maropoulos, S., Karagiannis, S., 2004. Rapid Prototyp. J. 10, 297. K. S. Kumar, T. E. Sparks and F. Liou, n.d. Department of Mechanical and Aerospace Engineering, Missouri University of Science and Technology, Rolla, MO. Letenneur, M., Brailovski, V., Kreitcberg, A., Paserin, V., Bailon-Poujol, I., 2017. J. Manuf. Mater. Proc. 1, 23. Michaeli, W., 2003. Extrusion Dies for Plastics and Rubber: Design and Engineering Computations. Hanser Verlag, Munchen. Middleman, S., 1995. Modeling Axisymmetric Flows Dynamics of Films, Jets, and Drops. Academic Press, San Diego, CA. Pandey, A., Pradhan, S.K., 2018. Mater. Today: Proc. 5, 12940. Ramanath, H.S., Chua, C.K., Leong, K.F., Shah, K.D., 2008. J. Mater. Sci. Mater. Med. 19, 2541. Thomas, J.P., Rodriguez, J.F., 2000. Modeling the fracture strength between fused deposition extruded roads. In: Solid Freeform Fabrication Proceedings. University of Texas at Austin, Austin, TX, pp. 16–23. Turner, B.N., Strong, R., Gold, S.A., 2014. Rapid Prototyp. J. 20, 192. Yardimci, M.A., Hattori, T., Guceri, S.I., Danforth, S.C., 1997. In: Bourell, D.L., Beaman, J.J., Crawford, R.H., Marcus, H.L., Barlow, J.W. (Eds.), Solid Freeform Fabrication Proceedings. University of Texas at Austin, Austin, TX.

Further reading Knyazeva, A., Travitzky, N., 2018. 3rd Inter. Conf. on Rheology and Modeling of Mater. IOP Publishing. IOP Conf. Ser.: J. Phys. Conf. Ser. 012020, 1045. Liao, Y.S., Chiu, Y.Y., 2001. Int. J. Prod. Res. 39, 3479. Melocchi, A., Parietti, F., Loreti, G., Maroni, A., Gazzaniga, A., Zema, L., 2015. J. Drug Delivery Sci. Technol. 30, 360. Noguera, R., Lejeune, M., Chartier, T., 2005. J. Eur. Ceram. Soc. 25, 2055. Zein, I., Hutmacher, D.W., Tan, K.C., Teoh, S.H., 2002. Biomaterials 23, 1169.

CHAPTER THREE

Testing: Comparison of AM data with traditionally fabricated

3.1 Tensile tests Although polymers represent one of the greatest markets of material produced by AM as various products (e.g., for biomedical devices), this chapter compares metallic and ceramic mechanical properties of samples obtained by AM and conventional fabrication method. The materials comprising this section are Ti alloys, Al alloys, and steel. However, before considering specific samples produced by AM it would be important to provide some relations of tensile properties. Assuming that a force, P, is acting normally on a small area, ΔA, of a test specimen allows writing for the stress, σ ¼ lim

dP

A!0 dA

¼

dp dA

(3.1)

dP ¼ σdA

(3.2)

Integration the expression yields

ð P ¼ σ dA σ¼

(3.3)

P A

(3.4)

The Hookean behavior of the stress-strain relation gives σ ¼ Ee

(3.5)

In Eq. (3.5), e is the average linear strain, which correlates the change in specimen dimension with its original length and which may be expressed as e¼

l  l0 Δl l ¼ ¼ 1 l0 l0 l0

Additive and Traditionally Manufactured Components https://doi.org/10.1016/B978-0-12-821918-8.00003-6

(3.6)

© 2020 Elsevier Inc. All rights reserved.

49

50

Additive and traditionally manufactured components

l0 is the original length within the gage length of the specimen, while Δl is the axial change resulting from the elastic deformation. Thus, Δl is often referred to as the “deformation.” Eq. (3.6), the linear strain, may also be expressed as ðl dl l  l0 e¼ ¼ (3.7) l0 l0 l 0 The strain in Eq. (3.6) is called the engineering strain and it is valid for small strains. A different and useful concept of defining the strain when deformation is considered in more practical terms is associated with the instantaneous change occurring in the specimen length while force is acting on it. Unlike in the engineering strain defined above where the reference was made to the constant gage length of the specimen in the present case we refer to the change in the linear dimension at each instant of the test. If dl is the amount by which the length l changes a strain can be defined similarly to Eq. (3.6) as ε¼

dl l

(3.8)

Integrating Eq. (3.8) an expression can be written as ðl2 ε¼

dl l2 ¼ ln l1 l

(3.9)

l1

The integration is performed between two arbitrary values of l, lengths l1 and l2. Clearly, when l is the instantaneous length and consideration is given to the original length, the integration limits should be changed resulting in ð li dl li (3.10) ε¼ ¼ ln l l 0 l0 ε is known as the natural, true or logarithmic strain. Eqs. (3.9) and (3.10) can be visualized basically as resulting from the summation of small changes between two values as given below:  X l1  l0 l2  l1 l3  l2 + + +⋯ (3.11) ε¼ l0 l1 l3 It is often required to switch between the two definitions of the strain. This can be performed easily. By using Eq. (3.7) it is possible to write l e ¼ 1 l0

(3.12)

Testing: Comparison of AM data with traditionally fabricated

51

or ðe + 1Þ ¼

l l0

(3.13)

and with Eq. (3.10) written as ε ¼ ln

l l0

(3.14)

one obtains ε ¼ ln ðe + 1Þ

(3.15)

It should be mentioned that for small strains 0.1 engineering and true strains are the same as seen below. Expand Eq. (3.15) in series ε ¼ ln ð1 + eÞ ¼ e 

e2 e3 + ⋯ 2 3

(3.16)

For small strains the terms where the exponents of e are larger than one can be neglected resulting in εffie

(3.17)

For more details on stress and the stress tensor, the reader is referred to the Mechanical Properties of Materials (Pelleg, 2012) and Mechanical Properties of Ceramics (Pelleg, 2014). In general, internal and external factors influence the performance of materials as indicated below: Internal factors of major influence are: (a) Grain size (b) Pores (c) Other flaws, such as microcracks and external factors are: (a) Composition (b) Specimen size (c) Specimen shape

3.1.1 Ti-6Al-4V: AM tensile properties Ti alloys are attractive for structural applications especially where lightweight components are required such as in aircraft, medical devices and implants, and many other industries because they have a high strength to weight ratio and good corrosion resistance. A very common Ti alloy is

52

Additive and traditionally manufactured components

the Ti-6Al-4V, which has been extensively studied in particular because of the mechanical properties and microstructure. In Fig. 3.1a and b the deposited Ti-6Al-4V microstructures obtained by DED while in Fig. 3.1c and d PBF AM techniques are shown. Due to the rapid solidification and the forming of the structure layerwise, the microstructure is anisotropic and consists of large prior-β grains that grow epitaxially across subsequent build layers. Common defects such as indicated in Fig. 3.2 are observed in the PBF process. The small, round gas entrapment pores are not detrimental to the ductility of parts made by AM, while lack of fusion porosity is harmful to the ductility of components. It is likely that the scatter observed in Fig. 3.3 are due to the defects in the structure.

Fig. 3.1 Micrographs of macrostructure and fine microstructure of Ti-6Al-4V fabricated by (a) directed energy deposition using a 2-kW laser power, with a zoom of inset shown in (b), and (c) powder-bed fusion using a 400-W laser with a zoom of inset shown in (d). In all micrographs, the build direction is vertical, and subsequent layers are horizontal. (Beese, A.M., Carroll, B.E., 2016. JOM 68, 724. Open access.)

53

Testing: Comparison of AM data with traditionally fabricated

Fig. 3.2 Micrographs from a powder-bed fusion specimen showing common defects in materials made by additive manufacturing: (a) round gas entrapment pores and (b) a sharp lack of fusion defect with partially melted powder particles. (Beese, A.M., Carroll, B.E., 2016. JOM 68, 724. Open access.) 1500 1350 Tensile strength (MPa)

1200 1050 900 750 600

DED as deposited

450

DED post treated

300

PBF as deposited

150

PBF post treated

0 0

2

4

6

8

10

12

14

16

18

20

Ductility (%)

Fig. 3.3 Tensile strength versus ductility for as-deposited and heat-treated samples fabricated by directed energy deposition and powder-bed fusion. (Data from Tables I, II, III, and IV. Beese, A.M., Carroll, B.E., 2016. JOM 68, 724. Open access.)

The tensile stress of specimens produced by DED and PBF in the asdeposited and posttreated conditions is shown in Fig. 3.3. The data are reproduced and presented as Tables 3.1–3.4. In the tables, the tensile strength and elongation and other pertinent data are listed for the various Ti-6Al-4V systems. Conventionally fabricated as cast Ti-6Al-4V has a yield strength of 895 MPa, UTS of 1000 MPa, and an elongation of 8%. The microstructure is lamellar with fine α colonies. Annealed Ti-6Al-4V for 1–4 h at 705–870°C,

Table 3.1 Mechanical properties of as-deposited Ti-6Al-4V fabricated by directed energy deposition. Laser

Mechanical properties Linear heat input (J/ mm)

Orientation

Tensile Elastic modulus yield (MPa) (GPa)

Tensile strength (MPa)

Ductility (%)

– – – – –

– – 950 976  24 950  2

1053  49 1035  26 – 1099  2 1025  10

– – 1 4.9  0.1 12  1

1063  2 5  1 1063  20 10.9  1.4 Zhang et al. (2009)

Type

Power (W)

Scan rate (mm/s)

Yb fiber

60–80

4

15–20

CW Nd:YAG 130–190 IΡG fiber 470 TRUMPF 1100–1200 DLD

8.5 16.7 12.5–14.2

15–22 28 77–96

Longitudinal Transverse – Longitudinal Longitudinal

IPG YLR– 12000

2000

10.6

189

Transverse Longitudinal

– –

950  2 960  26

CO2 laser CO2 laser

2400–2700 300

4–6 0.61

400–675 490

Transverse – Longitudinal

– – –

958  19 1064  26 14  1 1070 1140 6 1105  19 1163  22 4  1

0.01

33,000

Longitudinal

–

1005

1103

4

0.013

60,000

Longitudinal

–

990

1042

7

330 Optomec L850-R: low power 780 Optomec L850-R: high power

References

Yao et al. (2015) Zhang et al. (2001) Yu et al. (2012) Qiu et al. (2015)

Carroll et al. (2008) Palmer and Beese (2015) Zhai et al. (2015)

–

–

–

–

–

–

–

–

–

–

–

–

–

–

–

–

–

–

–

Optomec

–

–

–

Longitudinal, unmachined Transverse, unmachined Longitudinal, machined Transverse, machined –

–

–

Beese, A.M., Carroll, B.E., 2016. JOM 68, 724. Open access.

–

892  10

911  10

6.4  0.6

–

522  0

797  27

1.7  0.3

–

984  25

1069  19 5.4  1.0

–

958  14

1026  17 3.8  0.9

–

1069

1172

11

–

1077

973

11

Alcisto et al. (2011)

Keicher and Miller (1998) http://www. optomec.com/ 3d-PrintedMetals/lensMaterials/ (2015). Accessed 15 November 2015

Table 3.2 Mechanical properties of as-deposited Ti-6Al-4V fabricated by power bed fusion. Laser deposition

Type

ΜΤT 250 system Nd:YAG – ΜΤT SLM system SMYb:YAG fiber YAG

Nd:G Laser Cusing system

Mechanical properties

Orientation

Elastic modulus (GPa)

Tensile yield (MPa)

1

Longitudinal

–

125 600

0.76 0.27

175

710

0.25

– Longitudinal Transverse Longitudinal

94 105  5 102  7 –

250

1600

0.16

Longitudinal

109.2  3.1 1110  9

120– 200 – – –

–

–

Longitudinal

110  5

– – –

– – –

Longitudinal – Longitudinal 118  2.3 Transverse, 109.9 unmachined

Power Scan rate (W) (mm/s)

Linear heat input (J/mm)

200

200

95 160

Tensile strength (MPa)

Ductility (%)

References

910  9.9 1035  29

3.3  0.76

21

1125 1137  20 962  47 1166  6

29 47

1267  5

6 7.6  2 1.7  0.3 2.0  0.7 uniform 7.28  1.12

48

1095  10

8.1  3

25

8.2  0.3 6.5  0.6 11.9

26 33

990  5

1250 1206  8 1166  25 1321  6

1040  10 1140  10 1100  12 1211  31 736 1051

22

Yttrium fiber – Nd:YAG

– – – – 150– – 200

– – –

EOS M270 system EOS M270 system EOS

–

–

–

–

–

–

–

–

–

Beese, A.M., Carroll, B.E., 2016. JOM 68, 724. Open access.

Transverse, machined Transverse Longitudinal Longitudinal

112.4

986

1155

– –

1008 1080 1330 1400 1070  50 1250  50

1.6 4.4 5.5  1

32 31 23

Transverse Longitudinal Transverse Transverse

– – – 115

1050  40 1195  19 1143  30 1005

8.5  1.5 5  0.5 4.89  0.6 2.6

24

Longitudinal Transverse

110  10 110  10

1060  50 1230  50 1070  50 1200  50

10  2 11  3

60

1180  30 1269  9 1219  20 1190

10.9

59

Table 3.3 Mechanical properties of Ti-6Al-4V fabricated by directed energy deposition and subsequently heat treated. Mechanical properties

Posttreatment

Orientation

Elastic modulus (GPa)

700–730°C, 2 h

Longitudinal Transverse Longitudinal

116 112 –

1066 832 1000

1111 832 1073

5.2 0.8 9

Longitudinal

–

991

1044

10

Longitudinal Transverse – – Longitudinal, unmachined Transverse, unmachined Longitudinal, machined Transverse, machined

– – – – –

1052 1045 975  15 959  12 681  35

1153 1141 1053  18 1045  16 750  20

5.3 9.2 7.5  1 10.5  1 4.8  1.6

–

637  13

717  12

3.4  1.0

–

870  37

953  18

11.8  1.3

–

830  15

942  13

9.7  22

Low power, 760°C, 1 h, air cool High power, 760°C, 1 h, air cool 950°C, 1 h, quench; 538°C, 4 h, air cool 950°C, 1 h, air cool 950°C, 1 h, furnace cool 950°C, 1 h, furnace cool

Tensile yield (MPa)

Tensile strength (MPa)

Ductility (%)

References

Yadroitsev et al. (2014) Zhai et al. (2015)

Amsterdam and Kool (2009) Dinda et al. (2008) Alcisto et al. (2011)

Sub-β transverse anneal and age Anneal Mill anneal 1050°C, 1 h, air cool 1050°C, 1 h, furnace cool 900°C, 100 MPa, 2 h 920°C, 100 MPa, 4 h, furnace cool

–

–

839

900

12.3

Cottam and Brandt (2011)

– – – –

– – – –

827–965 958 931  16 900  14

896–1000 1027 1002  19 951  15

1–16 6.2 6.5  1 7.5  1

Griffith et al. (2000) Lewis and Schlienger (2000) Dinda et al. (2008)

Longitudinal Transverse –

118 114 –

949 899 850  2

1006 1002 920  1

13.1 11.8 17  2

Kobryn and Semiatin (2001)

Beese, A.M., Carroll, B.E., 2016. JOM 68, 724. Open access.

Qiu et al. (2015)

60

Table 3.4 Mechanical properties of Ti-6Al-4V fabricated by powder-bed fusion and subsequently heat treated. Mechanical properties Orientation

Elastic modulus (GPa)

Tensile yield (MPa)

Tensile strength (MPa)

100°C during fabrication 200°C during fabrication 540°C, 5 h, water quench 640°C, 4 h

Transverse Transverse Longitudinal Longitudinal

– – 112.6  30.2 –

– 1106  6 1118  39 1104  8

1314.9  15.6 – 1223  52 1225  4

Longitudinal –

1140  43

1214  24

Transverse

–

1152  11

1256  9

Transverse Transverse Longitudinal Longitudinal Transverse Transverse Transverse Longitudinal Transverse Longitudinal Longitudinal Transverse

– 117.4 114.6  2.2 101  4 110  29

850 1051 1026  35 965  16 900  101 962

700°C during fabrication 700°C, 1 h, 10°C/min cool 705°C, 3 h, air cool 730°C, 2 h, air cool 800°C, 2 h 800°C, 2 h, argon 800°C, 4 h, argon 850°C, 2 h, furnace cool 850°C, 5 h, furnace cool 900°C, 2 h; 700°C, 1 h,10°C/min cool 940°C, 1 h, air cool; 650°C, 2 h, air cool

116  10 114  10 114.7  3.6 112.0  3.4 118.8

860 min 860 min 955  6 909  24 908

940 1115 1082  34 1046  6 1000  53 1040 1228.1  32.4 930 min 930 min 1004  6 965  20 988

Longitudinal 115.6  2.4

899  27

948  27

Ductility (%)

References

4  1.2 11.4  0.4 5.36  2.02 7.4  1.6 uniform 3.2  2.0 uniform 39  1.2 uniform 6.5 11.3 9.04  2.03 9.5  1 1.9  0.8 5 8  1.5 10 min 10 min 1284  1.36 – 9.5

53 27 48 22

13.59  0.32 48

30 33 48 47 32 53 60 48 33

Additive and traditionally manufactured components

Posttreatment

1000°C, 1 h, furnace cool (4 h) 1000°C, 1 h, furnace cool (34 h) Heat treat variant 1 Heat treat variant 2 1020°C, 2 h, furnace cool 1050°C, 2 h in vacuum 1050°C, 2 h 1015°C, 0.5 h, air cool; 730°C, 2 h, air cool 1015°C, 0.5 h, air cool, 843°C, 2 h, furnace cool 1050°C, 1 h, water quench; 820°C, 2 h, air cool 900°C, 100 MPa, 2 h; 700°C, 1 h, 10°C/min cool 920°C, 100 MPa, 2 h 920°C, 100 MPa, 2 h 600–700°C, 2 h, furnace cool; 920°C, 103 MPa, 4 h, furnace cool 1050°C, 100 MPa, 2 h

Longitudinal 118  2.3 Longitudinal 103  11

960  19 944  8

1042  20 1036  30

13  0.6 8.5  1

112.8  2.9

925  14 826.87 804.77 835  5 870  15 760  19 – 798 822  25

1040  4 945.85 908.63 915  5 990  15 840  27 986.4  45.2 945 902  19

7.5  2 12.67 18.11 10.6  0.6 11.0  0.5 14.06  2.53 13.8  0.8 11.6 12.74  0.56

Longitudinal 1149  1.5

801  20

874  23

13.45  1.18

Longitudinal 96.7  5

913  7

1019  11

8.9  1

47

33

Transverse – – Longitudinal Longitudinal Longitudinal Transverse Transverse Longitudinal

98  3 – – – 117  1 114.7  09

Transverse Transverse

95  4 115.4

869  64 885

951  55 973

7.9  2 19

Transverse Transverse Longitudinal Transverse Transverse

– – – –

912 – 1000  60 930 –

1005 1088.5  26.3 1100  50 1020 1006.8  14.6

8.3 13.8  1.3 12.5  0.5 15.5  2 13.5  0.7

26 47

49 25 48 53 32 48

Testing: Comparison of AM data with traditionally fabricated

950°C, 0.5 h 950°C, 1 h, water quench; 700°C, 2 h, air cool

32 53 23 53

Beese, A.M., Carroll, B.E., 2016. JOM 68, 724. Open access.

61

62

Additive and traditionally manufactured components

Elongation

1200

25

1000

20

800

15

600

10

400

Wrought, annealed

Forged

Cast

EBM

EBM, HIP

DMLS

DMLS, HIP+HT

0

LENS, HT

0

LENS, HIP

5 DMD

200

Elongation (%)

YS

30

DMD, HIP+HT

Strength (MPa)

UTS 1400

Fig. 3.4 Tensile strength, yield strength, and elongation of Ti-6Al-4V alloy built using various AM processes (Pelleg, 2012). DMD, direct metal deposition; LENS, laser engineered net shaping; DMLS, direct metal laser sintering; EB, electron beam melting; HIP, hot isostatic pressing; HT, heat treatment. (Dutta, B., Froes, F.H., 2015. The additive manufacturing (AM) of titanium alloys, Chapter 24. Sci. Tech. Appl., 447. With kind permission of Elsevier and Dr. B. Dutta.)

air or furnace cooled has typically equiaxed microstructure with coarsened α colonies. The yield strength is 875 MPa, the UTS is 965 MPa, and the elongation is 13%. Tensile results of Ti-6Al-4V materials obtained by various AM techniques are compared in Fig. 3.4 with conventionally fabricated Ti-6Al-4V alloys in as-cast, forged and wrought-annealed conditions. Fig. 3.4 does not indicate dramatic changes in the tensile properties, however, the DMLS (direct metal laser sintering: AMF category of PBF) shows the greatest UTS and yield strength clearly on the expense of the elongation. The elongation is expected to increase after heat treatment as indeed observed (HIPped and heat-treated). The DMD (direct metal deposition) HIPped and heat-treated has as good elongation as the conventionally processed forged and wrought and annealed Ti-6Al-4V material. The major point of this comparison is that no deterioration of tensile properties occurred by any of the AM methods indicated in the figure and compared with the conventionally fabricated Ti alloys.

3.1.2 Al alloy AA6061: AM tensile properties One of the important aluminum alloys is the AA6061, which is heat-treatable and is composed of silicon and magnesium as its major alloying elements.

63

Testing: Comparison of AM data with traditionally fabricated

AA6061 is widely used in the aerospace and automobile industries due to its high strength to weight ratio, corrosion resistance, and ease of conventional fabrication by extrusion, rolling, and forging. It is also castable and readily joined by different methods, such as riveting, welding, brazing, soldering, and adhesive bonding. Usually, AA6061 alloy develops crack and porosity when using conventional manufacturing during solidification. One of the methods to get crack-free alloy is by preheating the AA6061 powder bed at relatively high temperatures resulting in nearly fully dense components using laser powderbed fusion (LPBF) AM technique. The powder morphology of spherical and semispherical particles is shown in Fig. 3.5. Oxidation and moisture pickup

(a) 40 35 Particle count

30 25 20 15 10 5 81-85

76-80

71-75

66-70

61-65

56-60

51-55

46-50

41-45

36-40

31-35

26-30

21-25

16-20

6-10

11-15

0-5

0

Powder Diameter (mm)

(b) Fig. 3.5 (a) SEM images of AA6061 gas atomized powder. Powder particles were spherical, and semispherical with satellites. (b) Histogram of AA6061 powder size distribution. (Uddin, S.Z., Murr, L.E., Terrazas, C.A., Morton, P., Roberson, D.A., Wicker, R.B., 2018. Addit. Manuf. 22, 405. With kind permission of Elsevier.)

64

Additive and traditionally manufactured components

has to be minimized because of the harmful effect of oxygen in promoting crack and porosity formation. During each build in the AM process using LBDF, the build chamber was continuously purged by Ar. The LPBF is also known as selective laser melting (SLM) or direct metal laser sintering (DMLS) and is a powder-bed fusion AM technology. A schematic representation of the system setup is seen in Fig. 3.6. The relative density of AA6061 cube coupons fabricated with different scan speeds and laser power was measured as part of the process parameter development study with and without powder-bed heating and is illustrated in Fig. 3.7. Bed heating was performed by using induction heating. The density measurements of three specimens in each condition were done using the relation of Eq. (3.18) D¼

A∙E AF

(3.18)

where A and F are measured weights of the specimen in air and water, respectively, and E is distilled water density. The nominal density of AA6061 is 2.7 g/cm3. Optical micrographs of AA6061 are shown in Fig. 3.8. They correspond to AA6061 produced without powder-bed heating. Large cracks are seen in both planes and also porosity is present. The former is attributed to large solidification range and the latter to entrapped gas. Cracks and porosity are eliminated when the process includes powderbed heating as illustrated in Fig. 3.9. Representative stress-strain diagrams are shown in Fig. 3.10. Table 3.5 lists the yield stress, UTS, and percent elongation and compares them with wrought AA6061 and wrought and heattreated (T6) AA6061. There is no dramatic difference in the results between them, in both with and without heat-treated AA6061. There is, however, a considerable difference in elongation as seen in Table 3.5. 3.1.2.1 Conventionally produced (AM) AA6061 Although in Table 3.5 AM produced AA6061 alloy is compared with conventionally fabricated annealed and wrought T6 alloy, a list of T6 alloys at various temperatures is presented in Table 3.6. However, let us recall that T6 Temper is a solution heat-treated and artificially aged Al alloy. It is commonly available in pretempered grades such as 6061-O (annealed), tempered grades such as 6061-T6. It is one of the most common alloys of aluminum for general-purpose use. AA 6061-T6 aluminum alloy is an Al-Mg-Si alloy that has wide acceptance in the fabrication of lightweight structures requiring high specific strength and good corrosion resistance. It is worth

65

Testing: Comparison of AM data with traditionally fabricated

1 Diverging optic 2 Focusing optic

1

3 Motorize deflection mirror

2

3 4 Laser beam 5 Curved focal plane if optics (1 and 2) were stationary 4 5 6

6 Flat focal plane since optics (1 and 2) are constantly adjusting

(a)

rake/re-coater rail system

x-y scanning system; mirrors and collimators

(i)

laser re-coater brush precursor powder

(ii)

inert (Ar) atmosphere

fabricated parts unused powder

supplier cylinder

build platform with induction heater

cylinder guide rods piston

(b) Fig. 3.6 (a) Scan-head of the AconityONE system consisting of the Galvo mirrors and f-q lens. The optical fiber receives the 1030 nm wavelength, near infrared laser from the laser module where Nd:YAG crystals were used for solid-state amplification. The scan-head is capable of rastering over an area of 400 mm in diameter. However, during heated bed configuration a reduced scan area with 200 mm in diameter is covered. (b) A schematic representation of the AconityONE LPBF system is used for the study. (i) Solid-state Nd:YAG laser of 1030 nm wavelength and 1 kW capacity and an induction heater (not shown in the figure) underneath the build platform that can raise the powder-bed brush is made of carbon fiber to impart flexibility and high heat resistance. (Uddin, S.Z., Murr, L.E., Terrazas, C.A., Morton, P., Roberson, D.A., Wicker, R.B., 2018. Addit. Manuf. 22, 405. With kind permission of Elsevier.)

66

Additive and traditionally manufactured components

Fig. 3.7 Relative density of AA6061 cube coupons fabricated using different laser power and scanning speed with and without heating of the powder bed. (Uddin, S.Z., Murr, L.E., Terrazas, C.A., Morton, P., Roberson, D.A., Wicker, R.B., 2018. Addit. Manuf. 22, 405. With kind permission of Elsevier.)

mentioning—for general knowledge—that quenching in water introduces residual stresses due to the presence of cooling gradient (between surface and center), and therefore the 6061-T651 often replaces T6 alloys. By introducing a 1%–3% stretching the residual stresses can be eliminated with a consequent possibility of machining and also avoiding distortion. T651 is also a solution heat treated, artificially aged alloy that has received a 1%–3% stretch which relieves residual stresses. The list is from the Clinton Aluminum of the Aluminum Association. In Table 3.7 the modulus of elasticity is shown. Tensile stress-strain relation can be seen in Fig. 3.11. The base material is the stress-strain curve of interest since the AA6061-T6 alloy is a high strength Al-Mg-Si alloy that contains manganese to increase ductility and toughness. In Table 3.8 the elongation is also presented in addition to the

Testing: Comparison of AM data with traditionally fabricated

67

Fig. 3.8 Microstructure of AA6061 specimens fabricated on the unheated powder bed. (a) and (b) illustrate the XY plane (perpendicular to build direction). (c) and (d) show the ZX plane (build direction). Cracks, porosity, melt-pool, and melt-track banding are evident in the microstructure. (Uddin, S.Z., Murr, L.E., Terrazas, C.A., Morton, P., Roberson, D.A., Wicker, R.B., 2018. Addit. Manuf. 22, 405. With kind permission of Elsevier.)

yield stress (σ x) and the rupture stress (σ t). Again the data of the base metal of 6061-T6 are included (see second raw).

3.1.3 Stainless steel 304L: AM tensile properties A popular AM steel is the AISI type 304L austenitic stainless steel (SS 304L). The tensile properties of this SS 304L made by direct energy deposition (DED) additive manufacturing from pre-alloyed stainless steel 304L powder are discussed in this section and compared with conventionally made tensile specimens. In Table 3.9 a summary of the mechanical properties of austenitic stainless steel is presented. Fig. 3.12 illustrates two walls and the base plate, the locations from which the tensile specimens were taken. Plots of uniaxial tension vs strain are seen in Fig. 3.13 for the specimens illustrated in Fig. 3.12. It seems (Fig. 3.13) that the baseplate specimen shows the highest stress. The tensile and yield strength

68

Additive and traditionally manufactured components

Fig. 3.9 Microstructure of AA6061 specimens fabricated on powder-bed heated to 500°C. (a) and (b) illustrate the XY plane (perpendicular to build direction). (c) and (d) show the ZX plane (build direction). Cracks, porosity, melt pool, and melt-track banding removed from the microstructure, and a columnar grain growth is observed in the build direction. (Uddin, S.Z., Murr, L.E., Terrazas, C.A., Morton, P., Roberson, D.A., Wicker, R.B., 2018. Addit. Manuf. 22, 405. With kind permission of Elsevier.)

variation from the baseplate distance are seen in Fig. 3.14. The specimens were obtained using laser-based direct energy deposition additive manufacturing. The yield strength, ultimate tensile strength, and ductility were higher in the lower linear heat input wall compared to the higher linear heat input wall as indicated in the figures. Table 3.10 summarizes the mechanical properties and the increase in ferrite numbers after the tensile test. The optical micrographs of broken longitudinal samples are illustrated in Fig. 3.15. Grain size is an important parameter affecting the mechanical properties and the kinetics of its growth can be expressed with Wang et al. (2016) as   dg k1 Q ¼ exp  dt 2g RT

(3.19)

69

Testing: Comparison of AM data with traditionally fabricated

Fig. 3.10 (a) Representative stress-strain diagram of LPBF fabricated AA6061 specimens as fabricated and T6 heat treated. (b) Solid cylinders built in X-direction and still on build plate; tensile specimens were machined out of these cylinders. (c) LPBF fabricated AA6061 tensile testing specimen after fractured. (Uddin, S.Z., Murr, L.E., Terrazas, C.A., Morton, P., Roberson, D.A., Wicker, R.B., 2018. Addit. Manuf. 22, 405. With kind permission of Elsevier.) Table 3.5 Comparison of mechanical properties of LPBF fabricated AA6061 and wrought annealed AA6061 for both annealed and T6 heat-treated conditions.

Heated powder bed-LPBF fabricated AA6061 Heated powder bed-LPBF fabricated AA6061 after T6 treatment Wrought AA6061-O Wrought T6 treated AA6061

Specimen Specimen Specimen Specimen

1 2 1 2

YS (MPa)

UTS (MPa)

Percent elongation at breakpoint

66 75 282 290

133 141 308 124

11 15 3.5 5.4

55 276

134 310

30 12

Uddin, S.Z., Murr, L.E., Terrazas, C.A., Morton, P., Roberson, D.A., Wicker, R.B., 2018. Addit. Manuf. 22, 405. With kind permission of Elsevier.

where g is the grain size, t is the time and k a kinetic constant. Clearly, Q is the activation energy for grain growth. The grain size is given as   Q g ¼ k1 ατexp  g02 RTP 2

(3.20)

Table 3.6 Physical properties of 6061-T6. Metric

English

Comments 3

2.70 g/cc

0.0975 lb/in

Mechanical properties

Metric

English

Comments

Hardness, Brinell

95

95

Hardness, Knoop

120

120

Hardness, Rockwell A

40

40

Hardness, Rockwell B

60

60

Hardness, Vickers

107

107

Tensile strength, ultimate

310 MPa 24.0 MPa @Temperature 32.0 MPa @Temperature 51.0 MPa @Temperature 131 MPa @Temperature 234 MPa @Temperature 290 MPa @Temperature 310 MPa @Temperature

45,000 psi 3480 psi @Temperature 4640 psi @Temperature 7400 psi @Temperature 19,000 psi @Temperature 33,900 psi @Temperature 42,100 psi @Temperature 45,000 psi @Temperature

AA; typical; 500 g load; 10 mm ball Converted from Brinell hardness value Converted from Brinell hardness value Converted from Brinell hardness value Converted from Brinell hardness value AA; Typical

371°C 316°C 260°C 204°C 149°C 100°C 24.0°C

AA; typical

700°F 601°F 500°F 399°F 300°F 212°F 75.2°F

Additive and traditionally manufactured components

Density

70

Physical properties

28.0°C 80.0°C 196°C 371°C 316°C 260°C 204°C 149°C 100°C 24.0°C 28.0°C 80.0°C 196°C

47,000 psi @Temperature 49,000 psi @Temperature 60,000 psi @Temperature 40,000 psi 1740 psi @Temperature 2760 psi @Temperature 4930 psi @Temperature 14,900 psi @Temperature 31,000 psi @Temperature 38,000 psi @Temperature 40,000 psi @Temperature 41,000 psi @Temperature 42,100 psi @Temperature 47,000 psi @Temperature

Testing: Comparison of AM data with traditionally fabricated

Tensile strength, yield

324 MPa @Temperature 338 MPa @Temperature 414 MPa @Temperature 276 MPa 12.0 MPa @Temperature 19.0 MPa @Temperature 34.0 MPa @Temperature 103 MPa @Temperature 214 MPa @Temperature 262 MPa @Temperature 276 MPa @Temperature 283 MPa @Temperature 290 MPa @Temperature 324 MPa @Temperature

18.4°F 112°F 321°F AA; typical 0.2% offset 700°F 0.2% offset 601°F 0.2% offset 500°F 0.2% offset 399°F 0.2% Offset 300°F 0.2% Offset 212°F 0.2% Offset 75.2°F 18.4°F 112°F 321°F

0.2% Offset 0.2% Offset 0.2% Offset 71

Continued

Elongation at break

17.0% ©Temperature 18.4°F 17.0% ©Temperature 75.2°F 18.0% @Temperature 112°F 18.0% @Temperature 212°F 20.0% @Temperature 300°F 22.0% @Temperature 321°F 28.0% @Temperature 399°F 60.0% @Temperature 500°F 85.0% @Temperature 601°F 95.0% @Temperature 700°F 12.0% @Thickness 0.0625 in. 17.0% @Diameter 0.500 in.

Comments

AA; typical AA; typical

Uddin, S.Z., Murr, L.E., Terrazas, C.A., Morton, P., Roberson, D.A., Wicker, R.B., 2018. Addit. Manuf. 22, 405. With kind permission of Elsevier.

Additive and traditionally manufactured components

17.0% @Temperature 28.0°C 17.0% @Temperature 24.0°C 18.0% @Temperature 80.0°C 18.0% @Temperature 100°C 20.0% @Temperature 149°C 22.0% @Temperature 196°C 28.0% @Temperature 204°C 60.0% @Temperature 260°C 85.0% @Temperature 316°C 95.0% @Temperature 371°C 12.0% @Thickness 1.59 mm 17.0% @Diameter 12.7 mm

English

72

Table 3.6 Physical properties of 6061-T6—cont’d Physical properties Metric

Table 3.7 Modulus of elasticity and other data.

Modulus of elasticity

68.9 GPa

10,000 ksi

AA; typical: average of tension and compression. Compression modulus is about 2% greater than tensile modulus 2.5 cm width  0.16-cm thick sidenotched specimen, Kt ¼ 17

Notched tensile strength Ultimate bearing strength Bearing yield strength Poisson’s ratio Fatigue strength

324 MPa

47,000 psi

607 MPa

88,000 psi

Edge distance/pin diameter ¼ 2.0

386 MPa

56,000 psi

Edge distance/pin diameter ¼ 2.0

0.330

0.330

Estimated from trends in similar Al alloys Completely reversed stress; 14,000 psi 96.5 MPa RR Moore machine/specimen @# of cycles @# of cycles 5.00e + 8 5.00e+8 Fracture 29.0 MPa- 26.4 ksi-in1/2 KIC; TL orientation toughness m1/2 Machinability 50% 50% 0–100 Scale of aluminum alloys Shear 26.0 GPa 3770 ksi Estimated from similar Al alloys modulus Shear 207 MPa 30,000 psi AA; typical strength Uddin, S.Z., Murr, L.E., Terrazas, C.A., Morton, P., Roberson, D.A., Wicker, R.B., 2018. Addit. Manuf. 22, 405. With kind permission of Elsevier.

Fig. 3.11 Tensile tests for welded material specimens. (Moreira, M.G.P., Santos, T., Tavares, S.M.O., Richter-Trummer, V., Vilac¸a, P., de Castro, P. M. S. T., 2009. Mater. Des. 30, 180. With kind permission of Elsevier.)

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Additive and traditionally manufactured components

Table 3.8 Material properties for FS welded specimens, data acquired in tensile tests. FSW

σ y (MPa) σ t (MPa) Elongation (%) Joint efficiency (%)

Base 6082-T6 Base 6061-T6 FSW 6082-T6 FSW 6082-T6 + 6061-T6

276.2 306.3 134.3 140.5

322.9 342.0 221.3 218.6

17.5 17.1 6.5 5.5

FSW 6061-T6

148.3

231.6

5.9

– – 68.5 67.7 (6082-T6) 63.9 (6061-T6) 64.2

FS stands for friction stir. Moreira, P.M.G.P., Santos, T., Tavares, S.M.O., Richter-Trummer, V., Vilac¸a, P., de Castro, P.M.S.T., 2009. Mater. Des. 30, 180. With kind permission of Elsevier.

here g0 is the initial grain size, TP is the peak temperature, and α and τ are defined as rffiffiffiffiffiffiffiffiffiffiffiffiffiffi 2πRTP (3.21) α¼ Q and τ¼

q=v 1 2πλe TP  T0

(3.22)

where q/v is the linear heat input, q is the laser power, v is the laser scanning speed, λ is the thermal conductivity, T0 is the preheat temperature, and τ is the time to preheat from T0 to TP. In the austenitic stainless steel under current discussion, the grains and ferrite dendrites are preferentially oriented (Fig. 3.15) along the highest thermal gradient and the microstructure of components are anisotropic, resulting in anisotropic mechanical properties. This is a consequence of the directional thermal gradients produced by the layer-by-layer build paths. To apply the kinetic grain growth model using Eqs. (3.19)–(3.22) calibration is required to evaluate the constants g20 and k1. Fig. 3.16 shows the average grain area, g2 versus τ, providing the constants g20 ¼ 1219 μm2 and k1 ¼ 1.9  109 μm2/s. Further one can apply the data for the Hall-Petch relation. Recall it relates the yield strength and the grain size according to k σ y ¼ σ 0 + pffiffiffi d

(3.23)

where σ y is the yield strength, d is the average grain diameter, and σ 0 and k are material constants. This relation is seen in Fig. 3.17.

Table 3.9 Summary of mechanical properties of AISI 304, 316, and 316L stainless steels fabricated by additive manufacturing compared with wrought properties reported in the literature. Stainless steel alloy

Laser power Scanning Linear heat (W) speed (mm/s) input (J/mm)

Density Orientation

Yield strength Tensile strength Elongation (MPa) (MPa) (%)

Directed energy deposition

Griffith et al. (2000)

304

–

–

–

100%

Longitudinal 448

710

59

Griffith et al. (1996, 2000)

316

–

–

–

100%

Transverse 324 Longitudinal 593

655 807

70 30

Xue et al. (2010) 316 [57] Zhang et al. 316 (2014) [20]

–

–

–

793 648–970

66 20–44

600–1400

2–10

75–500

Transverse 448 93.2%– Longitudinal 363–487 97.4% – Longitudinal 558

639

21

de Lima et al. (2014) [58] Yu et al. (2013) [2]

316

200–350

3–8

24–60

91%

Transverse Transverse

536 414–539

46 38–45

316L

570/750

13/17

45

99.6%

Longitudinal 490

685

51

Transverse Ma et al. (2013) [38]

316L

600–1650

7–23

69–90

96.5– 97.5%

280 400–440

580 430–510

62 14–20

304

205

520

40

316

220–270

520–680

40–45

352 207–261

Wrought

Guan et al. (2013) [6] Tolosa et al. (2010) [59]

–, Unspecified. Wang, Z., Palmer, T.A., Beese, A.M., 2016. Acta Mater. 110, 226. With kind permission of Elsevier.

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Additive and traditionally manufactured components

Fig. 3.12 Photograph showing the positions from which tensile specimens were extracted in each wall. x is the thickness direction, y is the longitudinal direction, and z is the build, or transverse, direction. (Wang, Z., Palmer, T.A., Beese, A.M., 2016. Acta Mater. 110, 226. With kind permission of Elsevier.)

800

Engineering stress (MPa)

700 600 500 400 Longitudinal: low power

300

Transverse: low power 200

Longitudinal: high power Transverse: high power

100

Baseplate 0 0

0.2

0.4

0.6

Engineering strain

Fig. 3.13 Representative engineering stress-strain curves of uniaxial tension samples extracted from the low power (2.3 kW) wall and high power (4 kW) wall in two directions, as well as a sample from the annealed baseplate. (Wang, Z., Palmer, T.A., Beese, A.M., 2016. Acta Mater. 110, 226. With kind permission of Elsevier.)

In summary, elongated grains grow along the build direction in the AM produced SS 304L resulting in anisotropic elongation (longitudinal specimen elongation < than transverse). The location-dependent yield strength can be described by the Hall-Petch relation and it is related to

77

Testing: Comparison of AM data with traditionally fabricated

(a)

(b) 680

Low power High power

380

Tensile strength (MPa)

Yield strength (MPa)

420

340 300 260 220

Low power High power

640

600

560

520

0 20 40 60 Distance from bottom of the wall (mm)

0 20 40 60 Distance from bottom of the wall (mm)

Fig. 3.14 Yield (a) and ultimate tensile strength (b) in longitudinal samples as a function of the distance of the sample gauge region from the baseplate. (Wang, Z., Palmer, T.A., Beese, A.M., 2016. Acta Mater. 110, 226. With kind permission of Elsevier.) Table 3.10 Summary of mechanical properties and increase in ferrite number after tensile tests. Low power wall

High power wall

Longitudinal Transverse Longitudinal Transverse Annealed n 5 11 n57 n 5 10 n58 baseplate n 5 4

Yield strength (MPa) Tensile strength (MPa) Elongation (%) Δ Ferrite number (FN)

337  29

314  6

277  27

274  7

265  9

609  18

606  13

581  20

560  12

722  14

48.2  2.5

56.4  5.8 41.8  3.5

50.5  6.7 62.3  2.6

0.9  0.3

1.3  0.4 0.9  0.3

1.0  0.3 41.7  4.1

Values in the table are average  standard deviation, where n indicates the number of samples tested in each condition. Wang, Z., Palmer, T.A., Beese, A.M., 2016. Acta Mater. 110, 226. With kind permission of Elsevier.

the location- and direction-dependent grain size indicating an anisotropic microstructure and anisotropy in the mechanical properties. The tensile strength and the elongation values of the samples of the DED AM produced 304L SS are lower than those obtained from the annealed base plate.

78

Additive and traditionally manufactured components

Fig. 3.15 Optical micrographs of broken longitudinal samples in which the build direction is vertical and subsequent build layers are horizontal in the images. Dashed lines indicate the transition between subsequent build layers. (a) Image of a sample extracted from the low power wall showing short grains within single layers. (b) Image of a sample extracted from the high power wall showing slightly elongated grains extending the full layers. (c) Zoom in of inset in (b) showing the lack of a sharp transition in microstructural features between subsequent build layers in which the bright phase is austenite and the dark features are skeletal d-ferrite dendrites. (d) Zoom in of inset in (c) showing δ-ferrite dendrites in the austenite matrix. (Wang, Z., Palmer, T.A., Beese, A.M., 2016. Acta Mater. 110, 226. With kind permission of Elsevier.)

3.1.3.1 Conventionally produced SS 304L Bulk and fiber 304L stainless steel were investigated for their mechanical properties. In Fig. 3.18a SEM micrograph of a fiber is shown. Images of fiber and bulk stainless steel, longitudinal and transverse are seen in Fig. 3.19. TEM micrographs of the fiber are shown in Fig. 3.20 indicating also the presence of some precipitate likely to be nanosized carbides forming during the fiber processing (Fig. 3.20b). Clearly, the formation of nanosized carbides during fiber processing can influence their mechanical properties. Engineering stress-strain curves of fiber and bulk of the 304L stainless steel tensile tests are plotted in Fig. 3.21. The tensile stress of the fiber is lower

79

Testing: Comparison of AM data with traditionally fabricated

9000 High power Low power g2 (mm2)

7000

5000 g2 = 11946t + 1219 3000 0.3

0.4

0.5

t (s)

0.6

Fig. 3.16 Average grain area versus time to heat for samples extracted 7 mm from the bottom and 15 mm from the top of both the low power wall and the high power wall, where average grain size was extracted from EBSD data (see original work). Data from the high power wall (black symbols) were used to fit the kinetic grain growth model, while data from the low power wall are shown in gray. The fitted line was used to calibrate the kinetic grain growth model to find the initial grain size, g0, and the kinetic constant, k1. This calibrated model was used to predict grain growth as a function of processing parameters for the low power wall. (Wang, Z., Palmer, T.A, Beese, A.M., 2016. Acta Mater. 110, 226. With kind permission of Elsevier.)

450 Yield strength sy (MPa)

sy = 695d -0.5 + 194 400 350 300 250 200 0.1

0.15

0.2

0.25

d -0.5 (mm-0.5)

Fig. 3.17 Yield strength versus d0.5, where d is the relevant grain diameter in the direction of the applied tensile load. The fitted line was used to determine the Hall-Petch parameters, σ 0 and k. (Wang, Z., Palmer, T.A, Beese, A.M., 2016. Acta Mater. 110, 226. With kind permission of Elsevier.)

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Additive and traditionally manufactured components

Fig. 3.18 SE-SEM micrograph of individual 304L stainless steel fibers: (a) longitudinal and (b) cross section. (Baldenebro-Lopez, F.J., Gomez-Esparza, C.D., Corral-Higuera, R., Arredondo-Rea, S.P., Pellegrini-Cervantes, M.J., Ledezma-Sillas, J.E., Martinez-Sanchez, R., Herrera-Ramirez, J.M., 2015. Materials 8, 451. Open access.)

Fig. 3.19 SE-SEM images of the AISI 304L samples: (a) cross; (b) longitudinal sections of fibers; (c) cross; (d) longitudinal sections of bulk material. (Baldenebro-Lopez, F.J., GomezEsparza, C.D., Corral-Higuera, R., Arredondo-Rea, S.P., Pellegrini-Cervantes, M.J., LedezmaSillas, J.E., Martinez-Sanchez, R., Herrera-Ramirez, J.M., 2015. Materials 8, 451. Open access.)

81

Testing: Comparison of AM data with traditionally fabricated

Fig. 3.20 TEM micrographs: (a) bright field; and (b) Z-contrast images of a fiber sample showing the presence of nanoscale precipitates. (Baldenebro-Lopez, F.J., Gomez-Esparza, C.D., Corral-Higuera, R., Arredondo-Rea, S.P., Pellegrini-Cervantes, M.J., Ledezma-Sillas, J.E., Martinez-Sanchez, R., Herrera-Ramirez, J.M., 2015. Materials 8, 451. Open access.)

600

Stress (MPa)

500 400

Bulk Fiber

300 200 100 0 0.0

0.1

0.2 0.3 0.4 Strain (mm/mm)

0.5

Fig. 3.21 Stress-strain curves of the bulk and fiber samples. (Baldenebro-Lopez, F.J., Gomez-Esparza, C.D., Corral-Higuera, R., Arredondo-Rea, S.P., Pellegrini-Cervantes, M.J., Ledezma-Sillas, J.E., Martinez-Sanchez, R., Herrera-Ramirez, J.M., 2015. Materials 8, 451. Open access.)

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Additive and traditionally manufactured components

Table 3.11 Mechanical properties of 304L stainless steel samples. HV Sample σ y (MPa)

Bulk Fiber

σ max (MPa)

ε (%)

E (GPa)

Long-section Cross-section

492  4.5 663.07  9.1 45.91  3.7 175.44  10.7 273.53  5.1 256.53  7.2 389  7.2 597.47  10.7 16.33  2.8 162.76  8.9 151.73  9.1 169.12  2.0

Baldenebro-Lopez, F.J., Gomez-Esparza, C.D., Corral-Higuera, R., Arredondo-Rea, S.P., PellegriniCervantes, M.J., Ledezma-Sillas, J.E., Martinez-Sanchez, R., Herrera-Ramirez, J.M., 2015. Materials 8, 451. Open access.

than that of the bulk. The plot of the bulk shows the classical tensile stress curve while the plot of the fiber indicates continued work hardening up to failure. The results of the tensile stress of the fibers and the bulk specimens are summarized in Table 3.11. The presence of porosities usually influences the mechanical properties of material. It turns out that porosity in the bulk was 2.13% while the fiber porosity was 5.84%. As indicated in Fig. 3.21 and Table 3.11 the fiber had overall lower properties than the bulk regardless if testing was performed on longitudinal or transverse specimens. Optical microscopy was used to determine the porosity on the polished cross sections. The pore diameters evaluated are 1.46  0.21 μm and 4.21  0.88 μm for bulk and fiber samples, respectively. The greater volume fraction of porosity and the size of porosities are thus detrimental to the mechanical properties. There is almost no difference in the chemical compositions between the fiber and the bulk (Table 3.12). (The main difference is in the Ni content.)

3.1.4 Ceramic In recent years, there has been a greater focus on the use of ceramic materials for AM than before. One of the reasons is that 3D printing allows the production of ceramic parts and shaping complex ceramic parts. However, the AM of ceramics did not reach yet the volume of production of polymers or Table 3.12 Chemical composition of the bulk and fiber specimens of 304L austenitic stainless steel (wt%). Sample

C

Si

Mu

Cr

Ni

P

S

Fe

Bulk Fiber

0.023 0.029

0.278 0.389

1.488 1.406

18.163 19.756

8.214 11.216

0.026 0.013

0.024 0.005

Balance Balance

Baldenebro-Lopez, F.J., Gomez-Esparza, C.D., Corral-Higuera, R., Arredondo-Rea, S.P., PellegriniCervantes, M.J., Ledezma-Sillas, J.E., Martinez-Sanchez, R., Herrera-Ramirez, J.M., 2015. Materials 8, 451. Open access.

Testing: Comparison of AM data with traditionally fabricated

83

metallic material. One of the reasons is probably the higher temperatures required in the case of ceramics. Continued research is going on to increase the commercial activity of ceramics produced by 3D printing. Further, it has to be noted, that many ceramic parts that come out of the 3D printer require to undergo a post-3D printing by an additional process of firing to change the fragile state of the “green object” and adding strength to it. The firing is performed at temperatures >800°C. Of the unlimited number of technical ceramics (oxides, carbides, and nitrides), the choice in this book is the consideration of alumina (an oxide), and the comparison of its mechanical properties with conventionally fabricated alumina. Note that we are talking on multistep or indirect AM process, but single-step powder-bed fusion by full melting is also a possibility for the ceramics AM fabrication. 3.1.4.1 AM alumina Here, a few general words on alumina is in place to indicate the technical importance of this ceramics: it has the formula Al2O3. Other special properties of this high-performance high purity oxide (99.7%) are: • High mechanical strength and wear resistance • Excellent thermal resistance • Outstanding electrical insulation characteristics • Extremely strong dielectric strength • Exceptional tribological properties • Can be processed with conventional abrasive tools • Low dissipation factor • High chemical resistance • Good metallizability • Very hard material structure Typical applications • Heat resistant ceramic components • Electrical performance insulators • Space applications • Dielectric for capacitors • Welding nozzles • Ceramic bushings • Wear-resistant components Porosity in alumina produced by AM has to be eliminated in order to get good mechanical properties in relatively high-density material. Extrusionbased AM process is capable of producing high-density alumina parts >95% of the theoretical density. The ceramic on-demand extrusion

84

Additive and traditionally manufactured components

Fig. 3.22 SEM image showing a typical microstructure of the Al2O3 produced via the CODE process. (Ghazanfari, A., Li, W., Leu, M., Watts, J., Hilmas, G., 2017. Int. J. Appl. Ceram. Technol. 14, 486. Open access.)

(CODE) process is a novel fabrication technique capable of making various and complex parts (even large size objects) with a near theoretical density of >98%. A typical microstructure of a printed alumina is shown in Fig. 3.22. The density variation for different stages of the process is in Table 3.13. The cross section is perpendicular to the printing direction and their grains are small equiaxed and measured by the linear intercept method. The grain size was evaluated by the relation X D ¼ 1:56 X

li

(3.24)

ni

Table 3.13 Amount of shrinkage and relative densities of parts at each stage. Size (mm)

Linear shrinkage (%)

Volumetric shrinkage (%)

As-printed 72.0  7.8  5.6 – – Dried 71.0  7.5  5.4 1.4  3.8  3.6 8.6 Sintered 62.8  6.3  4.6 12.8  19.2  17.9 42.1

Relative density (%)

57a 62a 98

a These densities are calculated by dividing mass of alumina powder by volume of the part. Ghazanfari, A., Li, W., Leu, M., Watts, J., Hilmas, G., 2017. Int. J. Appl. Ceram. Technol. 14, 486. Open access.

Testing: Comparison of AM data with traditionally fabricated

85

where D is the average grain size in μm, li is the length of each line in μm and ni is the number of intercept of each line. The cumulative distribution function for the Weibull distribution is     σ max m Pf ¼ 1  exp  (3.25) σ0 where Pf is the probability of failure, σ max is the maximum tensile stress in a test specimen at failure, σ 0 is the characteristic Weibull strength (corresponding to a Pf ¼ 0.632 or 63.2%), and m is Weibull modulus. Young’s modulus was calculated according to E¼

11Pl 3 768 Iδ

(3.26)

Young’s modulus is given as (N/m2), P is the load (N), l is the outer span of the fixture (m), I is the second moment of inertia and of the test specimen cross section about the neutral axis (m4) and δ is measured by the deflectometer and P is measured by a load-cell. The adjusted moment of inertia for a rectangular cross section with four chamfered edges of size c is given by ! bd 3 c 2 2 ð3d  2c Þ2 I¼  c + (3.27) 12 9 2 where b and d are width and height of the bar (m), respectively, and c is the chamfer size (m). Fig. 3.23 shows the Weibull plot of the flexural strength data. The mechanical properties are listed in Table 3.14 including the Weibull modulus (Fig. 3.24). Hardness data are also included in the table although it is discussed later in this chapter. A typical cross section showing a solid surface with no flaws is seen in a SEM micrograph in Fig. 3.25. Two fractured surfaces and the origin of the fracture are seen in Fig. 3.26. It is the opinion of the authors that the origin of the fracture are air bubbles or binder agglomerates. The flaws are found near the tensile surface. The Weibull characteristic strength is 385.3 MPa and the raw Weibull modulus is 8.3  0.943. Fracture surface and the surface of indentation for hardness measurements are illustrated in Fig. 3.24. In this method of AM fabrication (CODE) no flaws were observed. Calculation of the critical flaw size was performed using the Griffith criterion.   KIC 2 2c ¼ 2 (3.28) σf Y

86

Additive and traditionally manufactured components

99

Probability of failure, Pf (%)

90 80 50

20 10 5

1 250

300 350 400 Failure stress, sf (MPa)

450

500

Fig. 3.23 Weibull plot of the flexural strength data from Al2O3 test specimens. (Ghazanfari, A., Li, W., Leu, M., Watts, J., Hilmas, G., 2017. Int. J. Appl. Ceram. Technol. 14, 486. Open access.)

where 2c is the length of the flaw (m), KIC is the fracture toughness in MPa m0.5, σ f is the fracture stress (MPa), and Y is the stress intensity shape. The flaw location is near the tensile surface the place where σ f was evaluated. The values of Y are 1.77 and 1.13 for long and round flaws, respectively, and the estimated length, 2c, of the flaw is 102  34 μm for long flaws and 252  84 μm for the round flaws. The characteristic Weibull stress for flaws of volume originated and those of surface origin are given in the following equations:   1=mV  Li 1 V σV mV + 1 L0 2ðmV + 1Þ   ðσ 0 ÞV ¼ (3.29) 1 Γ +1 mV and   σ A L0 ðσ 0 ÞA ¼

 1=mA  d Li 1 +b mA + 1 L0 mA + 1 mA + 1   1 Γ +1 mA

(3.30)

Table 3.14 Mechanical properties of alumina parts produced by different additive manufacturing processes.

Process

CODE LCM (Schwentenwein and Homa, 2015) SLS Liu et al., 2007 RCb (Feilden et al., 2016) FEFb (Huang et al., 2009; Li et al., 2015a) 3DPb,c (Maleksaeedi et al., 2014) BJb (Gonzalez et al., 2016) a

Relative Young’s density modulus (GPa) (%)

Flexural strength (MPa)

Characteristic strength Weibull (MPa) modulus

Fracture toughness (MPa m0.5)

c

2c (μm) assuming long flaws

2c (μm) assuming round flaws

98 99

371  14 364  50 385.3 – 369a–383a –

8.3  0.943 4.5  0.1 11.2  0.955 –

19.8  0.6 102  34 – –

252  84 –

88 97 87–92

– 255  17 – – 236a–248a 297 327  20 219 –

– – 8.9  0.901 3.3  0.2 5.4  0.947 –

– – 18.6  0.8 89 14.4  0.9 –

– 218 –

85

–

–

–

–

–

–

–

–

54  14.5 Very lowd –

–

–

1.5  0.01 –

–

62

Original value converted to standard “B” bar using Eqs. (3.29) and (3.30) for fair comparison. Highest values in the paper are reported here. Vacuum infiltration was used to enhance the mechanical properties. d The compressive strength was only 132 MPa, so the flexural strength was minimal. Ghazanfari, A., Li, W., Leu, M., Watts, J., Hilmas, G., 2017. Int. J. Appl. Ceram. Technol. 14, 486. Open access. b

Hardness (GPa)

Fig. 3.24 Typical fracture surface (A) and indented surface (B). (Ghazanfari, A., Li, W., Leu, M., Watts, J., Hilmas, G., 2017. Int. J. Appl. Ceram. Technol. 14, 486. Open access.)

Fig. 3.25 A typical cross section under SEM showing a solid surface with no flaws. (Ghazanfari, A., Li, W., Leu, M., Watts, J., Hilmas, G., 2017. Int. J. Appl. Ceram. Technol. 14, 486. Open access.)

Fig. 3.26 Two typical fracture origins near the tensile surface of the Al2O3 flexure test specimens. (Ghazanfari, A., Li, W., Leu, M., Watts, J., Hilmas, G., 2017. Int. J. Appl. Ceram. Technol. 14, 486. Open access.)

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Testing: Comparison of AM data with traditionally fabricated

In the above equations, the subscripts stand for volume and surface area. In Eqs. (3.29) and (3.30) σ 0 is the Weibull material scale parameter, σ is the mean strength measured in experiments, Li and L0 are the lengths of inner and outer spans, respectively, m is the Weibull modulus, b and d are the width and height of the sample, respectively, V is the gage volume (or b  d  L0), and Γ is the gamma function. As known, the orientation of the specimen can affect the mechanical properties. It is expected that in the CODE process the effect is small. 3.1.4.2 Conventionally fabricated alumina Several grade aluminas were tested for the tensile properties investigating the influence of the strain rate. The properties of the aluminas before testing are listed in Table 3.15. In this table, ρ is the density and c is rate. AZR refers to 98% alumina, where the specimens were taken from 100 mm square tiles (it is reinforced with zirconia). The number of tests and the meaning of the devices and the mean tensile strength is shown in Tables 3.16 and 3.17. Note to Tables 3.16 and 3.17: A1 represents splitting tests at very low strain rate, A2 splitting tests at intermediate strain rate, B splitting tests in Hopkinson bar, and C spalling tests of long bars. Recall that in splitting Table 3.15 Properties of materials measured before testing. Material

ρ (kg/m3)

E (GPa)

c (m/s)

Grain size (μm)

A94 A98 A995 AZR

3658 3877 3905 4027

303 366 391 348

9108 9717 10,004 9292

8.3 2.4 10.4 2.0

Ga´lvez, F., Rodrı´guez, J., Sa´nchez Ga´lvez, V., 2000. J. Phys. IV France 10, Pr9-323. Free access.

Table 3.16 Number of tests done with each technique. Number of tests Testing device Material

A1

A2

C

D

94% Al2O3 98% Al2O3 99.5% Al2O3 AZR

5 6 10 7

4 0 5 1

9 6 8 6

8 5 7 5

Ga´lvez, F., Rodrı´guez, J., Sa´nchez Ga´lvez, V., 2000. J. Phys. IV France 10, Pr9-323. Free access.

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Additive and traditionally manufactured components

Table 3.17 Results of tests with each loading method. Mean tensile strength (MPa) Testing device Material

A1

94% Al2O3 98% A12O3 99.5% A12O3 AZR

161 179 161 155

(23) (21) (26) (12)

A2

B

181 (8) – 163 (29) 172

278 285 243 288

C

(28) (31) (43) (30)

358 329 271 322

(51) (78) (38) (35)

Ga´lvez, F., Rodrı´guez, J., Sa´nchez Ga´lvez, V., 2000. J. Phys. IV France 10, Pr9-323. Free access.

tensile stress a large compressive strength applied by the uniform load at one of the edges develops uniform tensile stress in the perpendicular direction. The compressive strength to failure should be at least three times more than the tensile strength to fail the material due to tension. It is expressed as σ f ¼ 2P=πLD

(3.31)

where σ f is tensile stress in the center fiber, P is the applied load, L is the length of the specimen, and D is the diameter of the cross section. The strain rate in the loading plane is expressed as   1 ∂σ t ε_splitting ¼ (3.32) E ∂t LoadPlane The spalling test at the fracture plane is given as   1 ∂σ t ε_spalling ¼ E ∂t FracturePlane

(3.33)

In Figs. 3.27–3.30 the standard tensile stresses and their deviations are illustrated. All samples tested show brittle fracture as seen in Fig. 3.31. No change in fracture mode was observed with variation of strain rates.

3.2 Compression tests Basically, in ductile metals, stress-strain relations are expected to show similar behavior in tension and compression. Failure occurs after necking when the ductile material is tested in tension, but in compression, barreling occurs. Not so, in brittle materials such as ceramics where the compressive strength is typically an order of magnitude greater than what it is in tension. For understanding the difference in tension and compression, let us assume

91

Testing: Comparison of AM data with traditionally fabricated

500 Material: A99

st (MPa)

400

300 C B 200 A

A 100

0 10-8

10-6

10-4

10-2

100

102

104

e¢eq (s ) -1

Fig. 3.27 Tensile strength and its standard deviation versus strain rate in 94% alumina. (Gálvez, F., Rodríguez, J., Sánchez Gálvez, V., 2000. J. Phys. IV France 10, Pr9-323. Free access.) 500 Material: A99

st (MPa)

400

300 C B 200 A

A 100

0 10-8

10-6

10-4

10-2

100

102

104

e¢eq (s ) -1

Fig. 3.28 Tensile strength and its standard deviation versus strain rate in 99.5% alumina. (Gálvez, F., Rodríguez, J., Sánchez Gálvez, V., 2000. J. Phys. IV France 10, Pr9-323. Free access.)

that a microcrack exists in the brittle material. When tension is applied, the crack propagates immediately because the plastic zone that exists at the crack-tip is too small—if at all. However, when compressive loading is applied, the crack is closed and the plastic zone in front of the crack-tip does not play a role in fracture. As a matter of fact, no crack-tip plastic zone exists

92

Additive and traditionally manufactured components

500 Material: A98 400

st (MPa)

C 300

200

B

A

100

0 10-8

10-6

10-4

10-2

100

102

104

e¢eq (s-1)

Fig. 3.29 Tensile strength and its standard deviation versus strain rate in 98% alumina. (Gálvez, F., Rodríguez, J., Sánchez Gálvez, V., 2000. J. Phys. IV France 10, Pr9-323. Free access.) 500 Material: AZR 400

st (MPa)

C 300

B

200 A

A 100

0 10-8

10-6

10-4

10-2

100

102

104

e¢eq (s ) -1

Fig. 3.30 Tensile strength and its standard deviation versus strain rate in alumina reinforced with zirconia. (Gálvez, F., Rodríguez, J., Sánchez Gálvez, V., 2000. J. Phys. IV France 10, Pr9-323. Free access.)

under compressive loading. The material can tolerate a high level of pressure without the occurrence of fracture. The strong bonds between atoms are responsible for high compressive strength of the brittle materials, especially it is true for ceramics. Since ductile metals are not as much influenced by

93

Testing: Comparison of AM data with traditionally fabricated

a)

b)

20 mm

c)

20 mm

d)

20 mm

10 mm

Fig. 3.31 (a) 1000 fracture surface of a splitting test of 94% alumina carried out in the static machine. Strain rate 106 s1. (b) 1000 fracture surface of a splitting test of 94% alumina carried out in Hopkinson bar. Strain rate 102 s1. (c) 1000 fracture surface of a spalling test of 94% alumina. Strain rate  103 s1. (d) 2000 fracture surface of a spalling test of 94% alumina. Strain rate 103 s1. (Gálvez, F., Rodríguez, J., Sánchez Gálvez, V., 2000. J. Phys. IV France 10, Pr9-323. Free access.)

such flaws, they should exhibit identical strengths in both tension and compression as indicated.

3.2.1 Ti-6Al-4V The Ti-6Al-4V alloy is commonly used in the aerospace industries, nuclear engineering, chemical industries, and as an implanted material due to its significant strength-to-weight ratio, resistance to corrosion and hightemperature creep. However, it is widely used as orthopedic implant due to the mentioned properties but first of most because of its biochemical compatibility. Thus, orthopedic and dental implants are favorable applications of the Ti-6Al-4V alloy provided the integration with the surrounding bone issue is solved. Micrograph obtained by optical microscope (OM) is seen in Figs. 3.32 and 3.33 and the heat treatments and other parameters are listed in

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Additive and traditionally manufactured components

Fig. 3.32 OM micrographs of the parallel section to the building platform (X-Y direction) of samples. (a) AB, (b) SR, (c) HT850, (d) HT950, and (e) HT1050 conditions. (Longhitano, G.A., Larosa, M.A., Jardini, A.L., de Carvalho Zavaglia, C.A., Ierardi, M.C.F., 2018. J. Mater. Process. Tech. 252, 202. With kind permission of J. Mater. Process. Tech.)

Table 3.18. DMLS in the table stands for direct metal laser sintering. In Fig. 3.33a the presence of elongated and oriented grains with the building direction is seen because heat removal occurs mainly by the base. In laser fusion processes the orientation of the grains of the solidified structure depends on the heat flow direction and the speed and scanning direction of the laser beam. Due to the high cooling rates from the β field temperature hexagonal martensitic transformations occur as illustrated in Fig. 3.34.

95

Testing: Comparison of AM data with traditionally fabricated

Fig. 3.33 OM micrographs of section parallel to the building platform (X-Y direction) of samples. (a) AB, (b) SR, (c) HT850, (d) HT950, and (e) HT1050 conditions. (Longhitano, G.A., Larosa, M.A., Jardini, A.L., de Carvalho Zavaglia, C.A., Ierardi, M.C.F., 2018. J. Mater. Process. Tech. 252, 202. With kind permission of J. Mater. Process. Tech.) Table 3.18 Parameters of heat treatments done in Ti-6Al-4V ELI produced by DMLS. Condition

Temperature (°C) Time (h) Healing rate (°C/min) Cooling

As-built (AB) Stress relieving (SR) HT850a HT950a HT1050a

– 650 850 950 1050

– 3 1 1 1

– 10 9.2 9.2 9.2

– Furnace Furnace Furnace Furnace

a Samples submitted to stress relieving before being removed from building platform. Longhitano, G.A., Larosa, M.A., Jardini, A.L., de Carvalho Zavaglia, C.A., Ierardi, M.C.F., 2018. J. Mater. Process. Tech. 252, 202. With kind permission of J. Mater. Process. Tech.

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Additive and traditionally manufactured components

Fig. 3.34 SEM micrographs of Ti-6Al-4V ELI produced by DMLS. (a) AB, (b) SR, (c) HT850, (d) HT950, and (e) HT1050 conditions. (Longhitano, G.A., Larosa, M.A., Jardini, A.L., de Carvalho Zavaglia, C.A., Ierardi, M.C.F., 2018. J. Mater. Process. Tech. 252, 202. With kind permission of J. Mater. Process. Tech.)

Compressive stress-strain curves are seen for the alloys heat-treated at different temperatures in Fig. 3.35. Table 3.19 summarizes mechanical properties obtained from the tests. Note that tensile test data are also included in Table 3.19. The fracture obtained by compression in a specimen is illustrated in Fig. 3.36. It is worth noting that barreling is observed before fracture indicating considerable elongation in the sample (see Table 3.19). As would be expected and as indicated in Section 3.2, the strength properties, yield, UTS, and even the elongation are greater in compression than in tensile tests.

97

Testing: Comparison of AM data with traditionally fabricated

2000 1750

Stress (MPa)

1500 1250 1000 750

AB SR HT850 HT950 HT1050

500 250 0 0

5

10

15 20 Strain (%)

25

30

35

Fig. 3.35 Stress-strain curves obtained by compression tests for different heat treatments conditions. (Longhitano, G.A., Larosa, M.A., Jardini, A.L., de Carvalho Zavaglia, C.A., Ierardi, M.C.F., 2018. J. Mater. Process. Tech. 252, 202. With kind permission of J. Mater. Process. Tech.)

3.2.2 Conventionally fabricated Ti-6Al-4V Compression tests were performed on Ti-6Al-4V at various strain rates and various temperatures. The flow stresses are shown in Fig. 3.37 at 25°C and 500°C at the strain rate indicated. The alloys tested at all temperatures at a strain rate of 1 s1 are seen in Fig. 3.38. The 0.1 s1 strain rate effect at all temperatures is seen in Fig. 3.39. Table 3.20 summarizes the results of the strain rate variations at the test temperatures obtained by compression. It is observed that the flow stress increases with increasing the strain rate. The temperature has an opposite effect, namely, the flow stress decreases with increasing temperature. The forms of the compression tests are shown in Fig. 3.40. Specimens compressed at uniaxial compression at strain rates above 0.01 s1 show shear bands up to 500°C after specimen height reduction of 70%. However, below 30% height reduction the specimens deformed uniformly except at 500°C. The overall compression tests were carried out at constant strain rates of 0.01, 0.1, and 1 s1 (see Fig. 3.38a–c) to height-reduction of 30%, 50%, and 70% in the temperature range of 25–500°C (namely, 25°C, 100°C, 200°C, 300°C, and 500°C). The flow stress data are analyzed in terms of strain rate and temperature sensitivities. The flow rate sensitivity is calculated by     σ1 ε1 = ln (3.34) m ¼ ln σ2 ε2

98

Table 3.19 Mechanical properties obtained by tensile and compression tests for different heat treatments conditions in Ti-6Al-4V ELI produced by DMLS. Tensile test

Compression test

σ y0.2% (MPa)

σ u (MPa)

ε (%)

E (GPa)

σ y0.2% (MPa)

σ u (MPa)

ε (%)

E (GPa)

AB SR HT850 HT950 HT1050 Commercial Ti-6Al-4Va

1015 + 10 1040  7 945 + 6 860  5 787  4 962

1218 + 2 1101  5 1003  4 926  3 869  3 1045

5.9 + 1.0 7.8  0.7 8.1  0.3 10.5  0.6 11.5  1.0 15

109 + 4 111  1 114 + 1 114  2 114  1 106

1134  9 1237  47 1078  4 954  8 851  7 –

1920 + 58 1950  34 1961 + 39 1943  18 1797  18 –

24.3 + 1.1 28.6  0.6 31.9  0.4 34.1  0.8 33.7  0.4 –

118 + 2 118  3 119 + 3 120  3 118  4 –

a Annealed commercial Ti-6Al-4V alloy results from Larosa et al. (2015) work. Longhitano, G.A., Larosa, M.A., Jardini, A.L., de Carvalho Zavaglia, C.A., Ierardi, M.C.F., 2018. J. Mater. Process. Tech. 252, 202. With kind permission of J. Mater. Process. Tech.

Additive and traditionally manufactured components

Condition

Testing: Comparison of AM data with traditionally fabricated

99

Fig. 3.36 Ti-6Al-4V ELI produced by DMLS sample fractured by compression test: fracture in 45 degrees to the load axis. (Longhitano, G.A., Larosa, M.A., Jardini, A.L., de Carvalho Zavaglia, C.A., Ierardi, M.C.F., 2018. J. Mater. Process. Tech. 252, 202. With kind permission of J. Mater. Process. Tech.)

Compression tests above 500°C are illustrated in Fig. 3.41. Corrected and uncorrected flow curves are indicated at three temperatures at 0.001, 0.1, and 10 s1 strain rates. The strain rate sensitivity, m is illustrated in Fig. 3.42. The B microstructure together with A microstructure are seen in Fig. 3.43. The grain size of A is 100 μm and that of B is 400 μm. The strain rate sensitivity, m is shown in Table 3.21. Flow softening in the flow curves is observed and the data are listed in Table 3.22. Note that two kinds of softening are observed: heating related softening and microstructure related softening. The stress exponent was found to be n ¼ 4  0.7 over the entire strain rate-temperature range. Analysis of the hot deformation results indicates that dislocation glide process controls the flow stress as evidenced by the fact that the flow curves are identical despite the different grain sizes and by the large values of the stress exponent n 4.

3.2.3 Al alloys-Al 60613 Only limited information is available in the literature on the compressive strength of additively manufactured Al 6061, and knowledge on the fundamental metallurgy of AM is lacking. Below recent information on AM of AA 6061 is presented (published in a dissertation in Norway).

100

Additive and traditionally manufactured components

True Strain - True Stress curve of Ti-6 Al-4V Alloy at 25°C

Flow Stress (Mpa)

1200 1100 1000 900

1 s-1 0.1 s-1 0.01 s-1

800 700 600 0.1

0.2

(a)

0.3

0.4

0.5

0.6

True Strain

True Strain-True Stress curve of Ti-6Al-4V alloy at 500°C 800 Flow Stress (Mpa)

750 700 650 600

1 s-1 0.1 s-1 0.01 s-1

550 500 450 400 0.1

(b)

0.2

0.3

0.4

0.5

0.6

True Strain

Fig. 3.37 True strain-true stress curves of Ti-6Al-4V alloy at (a) 25°C and (b) at 500°C. (Arulselvan, M., Ganesan, G., 2013. Int. J. Recent Tech. Eng. (IJRTE) 2, 47. Open access.)

SEM images of the powders are seen in Fig. 3.44. In many instances, smaller particles are attached to larger ones, commonly known as satellites. Optical micrographs of polished samples are seen in Fig. 3.45. In Fig. 3.46 binary images of the powder are shown in above (top) view. The side view can be found in the work of Rønneberg (2016). Micropores disrupted and large oxides and cracks are seen in the images. The powder is used to build a component in PBF process, which is one of the AM techniques. Following characterization of the powder mechanical properties evaluation was performed in both tensile and compressive testing. It seems that the 6061 alloy is weaker than the AlSi10Mg one (not discussed in this book), probably because of the large number of cracks found in the 6061 alloy. This can be seen in Fig. 3.47.

101

Testing: Comparison of AM data with traditionally fabricated

Flow Stress(MPa)

True Strain - True Stress Curve of Ti-6Al-4V alloy at 1S-1 or all of the temperatures tested 1200 1100 1000 900 800 700 600 500 400

(a)

25 o C 100 o C 300 o C 500 o C 0.1

0.2

0.3

0.4

0.5

0.6

True Strain

Flow Stress(Mpa)

True Strain-True Stress curve of Ti-6Al-4V alloy at 0.1 S-1 for all the temperature tested 1200 1100 1000 900 800 700 600 500 400

25 o C 100 o C 300 o C 500 o C 0.1

0.2

0.3

0.4

0.5

0.6

(b)

True Strain

Flow Stress(MPa)

True Strain-True Stress curve of Ti-6Al-4V alloy at 0.01 S-1 for all the temperature tested 1400 1200 1000 800 600 400

25 o C 100 o C 300 o C 500 o C 0.1

0.2

0.3

0.4

True Strain

0.5

0.6

(c)

Fig. 3.38 True strain-true stress curves of Ti-6Al-4V alloy for all the temperatures tested at the strain rates indicated. (a) 1 s1, (b) 0.1 s1, and (c) 0.01 s1. (Arulselvan, M., Ganesan, G., 2013. Int. J. Recent Tech. Eng. (IJRTE) 2, 47. Open access.)

The consolidated components are characterized by the compressive stress-strain relation illustrated in Fig. 3.48. The specimens after compression are seen in Fig. 3.49. In the figure the Cs are:

102

Additive and traditionally manufactured components

Flow Stress(Mpa)

Effect of the strain rate on the flow stress at a plastic strain of 0.1 for all of the temperatures tested 1200 1100 1000 900 800 700 600 500 400

25°C 100°C 300°C 500°C 0.1

0.2

0.3 -1

Strain Rate S

Fig. 3.39 Effect of the strain rate on the flow stress at a plastic train of 0.1 for all of the temperatures tested. (Arulselvan, M., Ganesan, G., 2013. Int. J. Recent Tech. Eng. (IJRTE) 2, 47. Open access.) Table 3.20 Strain rate sensitivity values calculated for different temperatures, strains and strain rate ranges. Strain Strain rate

ε 5 0.1

ε 5 0.2

ε 5 0.3

0.006 0.016

0.009 0.019

0.19 0.014

0.001 0.028

0.001 0.019

0.005 0.010

0.015 0.004

0.013 0.021

0.014 0.024

0.027 0.019

0.027 0.012

0.024 0.007

25°C

0.01–0.1 0.1–1.0 100°C

0.01–0.1 0.1–1.0 300°C

0.01–0.1 0.1–1.0 500°C

0.01–0.1 0.1–1.0

Arulselvan, M., Ganesan, G., 2013. Int. J. Recent Tech. Eng. 2, 47. Open access.

C0 Cylinder with length direction normal to the fabrication plane. C45 Cylinder with length direction 45 relative to the fabrication plane and C90 Cylinder with length direction in the fabrication plane.

Testing: Comparison of AM data with traditionally fabricated

103

Fig. 3.40 Different forms of the compression test results. (Arulselvan, M., Ganesan, G., 2013. Int. J. Recent Tech. Eng. (IJRTE) 2, 47. Open access.)

3.2.4 Conventionally fabricated Al 6061 The fabrication of the Al 6061 alloy was performed by near-liquidus casting. The cylindrical samples for the testing were cut from billets. The fabrication is known as semiliquid processing. The experimental scheme is illustrated in Fig. 3.50. Heating of the specimens was done in an electric resistance furnace. Graphite slices were placed between the specimen and the compression head to reduce the effect of the friction force. The semisolid deformation temperature was in the range 582–652°C and the heating rates are seen in the figure (temperatures: 585°C, 595°C, and 605°C) with a 10 s hold time at the high temperature. The semisolid compression was performed at strain rates of 0.01, 0.1, 1, and 10 s1, respectively. The microstructural changes were followed by optical microscopy. The microstructure before the compression tests is illustrated in Fig. 3.51. The microstructures of semisolid samples after compression at different strain rates and different temperatures are illustrated in Figs. 3.52–3.53. In Fig. 3.54 the microstructures at 605°C for various strain rates (indicated in the figure) are illustrated. These microstructures correspond to the compression stress-strain curves presented in Fig. 3.55. As expected the peak stress illustrated in Fig. 3.55 increases with decreasing the deformation temperature. The peak stress also increases with increasing strain rate. The change in the flow stress beyond the peak stress is not affected very much (see, e.g., Fig. 3.55d).

104

Additive and traditionally manufactured components

Fig. 3.41 Measured and temperature-corrected flow curves for Ti-6Al-4V with the B microstructure upset at nominal test temperatures of (a) 815°C; (b) 900°C; or (c) 955°C. (Semiatin, S.L., Seetharaman, V., Weiss, I., 1999. Mater. Sci. Eng. A263, 257. With kind permission of Elsevier.)

105

Testing: Comparison of AM data with traditionally fabricated

Fig. 3.42 Dependence of the strain rate sensitivity (m value) on strain determined in strain-rate-change tests on Ti-6Al-4V with the B microstructure. (Semiatin, S.L., Seetharaman, V., Weiss, I., 1999. Mater. Sci. Eng. A263, 257. With kind permission of Elsevier.)

Fig. 3.43 Optical micrographs of Ti-6Al-4V program material: (a) Condition A and (b) condition B. (Semiatin, S.L., Seetharaman, V., Weiss, I., 1999. Mater. Sci. Eng. A263, 257. With kind permission of Elsevier.) Table 3.21 Strain rate sensitivity data determined from peak stresses of continuous flow curves.a Temperature

m (εp) at ε_ ¼ 1023–1021 s21

1021–101 s21

815 900 955

0.148 0.183 0.218

0.058 0.101 0.105

a Ti-6Al-4V with B (colony) microstructure. Semiatin, S.L., Seetharaman, V., Weiss, I., 1999. Mater. Sci. Eng. A263, 257. With kind permission of Elsevier.

106

Additive and traditionally manufactured components

Table 3.22 Flow softening behavior.a Temperature Strain rate (°C) (s21)

815 815 815 900 900 900 955 955 955

103 101 101 103 101 101 103 101 101

Δσ/σ p at ε 50.50 Heating related softening

Microstructure related Total softening softening

0.00 0.09 0.16 0.00 0.09 0.15 0.00 0.10 0.17

0.26 0.29 0.12 0.23 0.31 0.13 0.16 0.23 0.05

0.26 0.38 0.28 0.23 0.40 0.28 0.16 0.33 0.22

a

Ti-6Al-4V with B (colony) microstructure. Semiatin, S.L., Seetharaman, V., Weiss, I., 1999. Mater. Sci. Eng. A263, 257. With kind permission of Elsevier.

100 mm

EHT = 15.00 kV Signal A = SE2 WD = 10.1 mm Mag = 150 X

10 mm

EHT = 15.00 kV Signal A = SE2 WD = 10.1 mm Mag = 500 X

Fig. 3.44 SEM images of powders. (Left) AA6061 low magnification. and (right) AA6061 high magnification. (Rønneberg, T., 2016. Characterization of Aluminium Components Produced by Additive Manufacturing. Norwegian University of Science and Technology, Trondheim 2016. With kind permission of Dr. Tobias Rønneberg.)

3.2.5 AM stainless steel 304L One of the alloys used in the additive manufacturing (AM) process is the 304L stainless steel (SS). Differences between the AM fabricated materials as compared to wrought materials might be expected, due to possible differences in porosity (voids) grain size, and residual stress levels. In the next

Testing: Comparison of AM data with traditionally fabricated

107

Fig. 3.45 Overview images of polished specimens. (Left) AA6061 above view and (right) AA6061 side view. (Rønneberg, T., 2016. Characterization of Aluminium Components Produced by Additive Manufacturing. Norwegian University of Science and Technology, Trondheim 2016. With kind permission of Dr. Tobias Rønneberg.)

section, conventionally fabricated 304L steel will be discussed. In this section, the effect of strain rate on the compressive stress is considered and compared with wrought 304L test under similar conditions. Compressive and tensile stress-strain curves are compared in Fig. 3.56. Strain rates of 500, 1500, and 2300 s1 were used to evaluate the dynamic stress-strain curves. Wrought 304L SS tests were also performed for comparison with similar chemical composition and the same strain rate. The microstructures of the AM 304L SS and the wrought steel are compared in Fig. 3.57. The stress at the front face, σ 1 and the stress at the back face, σ 2 were calculated with the following respective equations: σ1 ¼

A0 E0 ðεi + εr Þ As

(3.35)

A0 E0 εt As

(3.36)

σ2 ¼

In Eqs. (3.35)–(3.36) the strains, εi, εr, and εt are the incident, reflected, and transmitted bar engineering strains, respectively. E0 is Young’s modulus of the incident/transmission Kolsky compression bar material A0 and As are the cross-sectional areas of the bars and specimens, respectively. The test bars and their orientation is seen in Fig. 3.58. As seen in the figure the laser beam was aligned in the Z-direction of the specimen while the crosshatches were

108

Additive and traditionally manufactured components

Fig. 3.46 Binary images of AA6061 above view. (a) Micrograph. (b) Metallurgical pores. (c) Disrupted oxides. (d) Cracks. (e) Large oxides. (f) Overlay image. (Rønneberg, T., 2016. Characterization of Aluminium Components Produced by Additive Manufacturing. Norwegian University of Science and Technology, Trondheim 2016. With kind permission of Dr. Tobias Rønneberg.)

Testing: Comparison of AM data with traditionally fabricated

109

Fig. 3.47 SEM images of cracks in AA6061 samples. (Rønneberg, T., 2016. Characterization of Aluminium Components Produced by Additive Manufacturing. Norwegian University of Science and Technology, Trondheim. With kind permission of Dr. Tobias Rønneberg.)

110

Additive and traditionally manufactured components

250

Stress [MPa]

200

150 C0 C45

100

C90 50

0 0

10

20

30

40

50

60

70

80

90

Strain [%]

Fig. 3.48 Compressive stress-strain plot for AA6061. (Rønneberg, T., 2016. Characterization of Aluminium Components Produced by Additive Manufacturing. Norwegian University of Science and Technology, Trondheim. With kind permission of Dr. Tobias Rønneberg.)

Fig. 3.49 AA6061 specimen after compression testing. (a) C0, (b) C45, and (c) C90. (Rønneberg, T., 2016. Characterization of Aluminium Components Produced by Additive Manufacturing. Norwegian University of Science and Technology, Trondheim. With kind permission of Dr. Tobias Rønneberg.)

oriented in the X- and Y-directions. The lengths of the bars were in the Z-direction (Z bar) and X-direction (X bar). In Fig. 3.59, a comparison of the engineering stresses in both ends of the specimens is seen. The stresses at 3.35 and 3.36 represent the front and back faces, respectively. The engineering strain rates and the engineering strain in the specimens are evaluated by 2C0 ε_ ¼  εr (3.37) Ls ð 2C0 t εr dt (3.38) ε¼ Ls 0

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Testing: Comparison of AM data with traditionally fabricated

Temperature

Compression Holding time 10 s

1 °C/s Cooling

5 °C/s

Time

Fig. 3.50 Experimental scheme of 6061Al alloy. (Shang, S.-Z., Lu, G.-M., Tang, X.-L., Zhao, Z.-X., Wu, C.-M., 2010. Trans. Nonferrous Met. Soc. China 20, 1725. With kind permission of Elsevier.)

(a)

(b)

250 mm

50 mm

Fig. 3.51 Microstructures of 6061Al alloy before compression process: (a) conventional as-cast sample and (b) semisolid sample fabricated by near-liquidus casting. (Shang, S.-Z., Lu, G.-M., Tang, X.-L., Zhao, Z.-X., Wu, C.-M., 2010. Trans. Nonferrous Met. Soc. China 20, 1725. With kind permission of Elsevier.)

where C0 is the elastic stress wave speed in the bar material and Ls is the gage length of the specimen. Note that a Kolsky bar is an apparatus for testing the dynamic stress-strain response of materials. A schematic illustration of the Klosky compression bar is seen in Fig. 3.60. Further to Fig. 3.56, a comparison between the dynamic compressive stress-strain curves at various strain rates of the wrought and the AM

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Additive and traditionally manufactured components

Fig. 3.52 Microstructures of semisolid 6061Al alloy after compression process at 585°C: (a) 0.01 s1; and (b) 0.1 s1. (Shang, S.-Z., Lu, G.-M., Tang, X.-L., Zhao, Z.-X., Wu, C.-M., 2010. Trans. Nonferrous Met. Soc. China 20, 1725. With kind permission of Elsevier.)

Fig. 3.53 Microstructures of semisolid 6061Al alloy after compression process at 595°C: (a) 0.01 s1 and (b) 10 s1. (Shang, S.-Z., Lu, G.-M., Tang, X.-L., Zhao, Z.-X., Wu, C.-M., 2010. Trans. Nonferrous Met. Soc. China 20, 1725. With kind permission of Elsevier.)

304L specimen in the Z direction are seen in Fig. 3.61. As can be seen, the two groups are assembled, each in its own fabrication method. The elasticplastic characteristics of the two groups are quite similar. The thick lines represent longitudinal orientation and the thin lines transverse specimen orientations (directions). The AM SS shows in the Z-direction higher yield and flow stress than the wrought steel (by about 20%). However, at large strains, the trend seems to be reversed and the wrought steel indicates a higher stress level. Annealed (along the X-direction) and as-deposited (along the Z-direction) 304L SS are compared in Fig. 3.62. As would be expected the annealed samples exhibit lower yield and flow stress. Results of

Testing: Comparison of AM data with traditionally fabricated

113

Fig. 3.54 Microstructures of semisolid 6061Al alloy after compression process at 605°C: (a) 0.01 s1; (b) 0.1 s1; (c) 1 s1; and (d) 10 s1. (Shang, S.-Z., Lu, G.-M., Tang, X.-L., Zhao, Z.-X., Wu, C.-M., 2010. Trans. Nonferrous Met. Soc. China 20, 1725. With kind permission of Elsevier.)

compressive stress at two locations, the center and edge of the AM 304L SS are shown in Fig. 3.63. In Table 3.23 details of the testing conditions, processing, and the results of the measurements are summarized. Further comparison between AM 304L and wrought SS is seen in Fig. 3.64. Notice in the figure that the quasi-static constitutive yield strength of the AM 304L SS is much higher than the wrought material. Further, the as-built AM 304L SS is much higher than the recrystallized samples (by 60%). All three materials are shown in the figure exhibit similar work hardening rate. In the inset of the figure strain rate, jumps were performed (at 233 K) which are invariant in the three materials indicating that the increase in strength of the AM as-built material is a consequence of its dendritic substructure and not a result of different defect (dislocation) content or of residual stresses. Optical micrographs show these materials in the fabrication conditions indicated in Fig. 3.64.

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Additive and traditionally manufactured components

–10

(a)

0.01 s–1 0.1 s–1 1 s–1 10 s–1

True stress, s /MPa

–10 –8 –6 –4 –2 0

–0.2

–0.4 –0.6 True strain, e

–4 –2

0

–2

–0.2

–0.4 –0.6 True strain, e

–0.8

–0.2

–0.4 –0.6 True strain, e

(d)

–1.0

–0.8

–1.0

585 ˚C 595 ˚C 605 ˚C

–16 True stress, s /MPa

–4

0

–1.0

585 ˚C 595 ˚C 605 ˚C

–6

0.01 s–1 0.1 s–1 1 s–1 10 s–1

–6

–20

(c) –8 True stress, s /MPa

–0.8

(b)

–8 True stress, s /MPa

–12

–12 –8 –4

0

–0.2

–0.4 –0.6 True strain, e

–0.8

–1.0

Fig. 3.55 True stress-true strain curves of semisolid 6061Al alloy during compression process at different temperatures and strain rates: (a) 595°C; (b) 605°C; (c) 0.01 s1; and (d) 10 s1. (Shang, S.-Z., Lu, G.-M., Tang, X.-L., Zhao, Z.-X., Wu, C.-M., 2010. Trans. Nonferrous Met. Soc. China 20, 1725. With kind permission of Elsevier.)

3.2.5.1 Conventionally fabricated stainless steel 304L Compressive and tensile stresses of 304L SS produced by a method known as equal channel angular pressing (ECAP) at high temperatures (of 700°C), which is a technique for producing ultrafine-grained materials (grain size in the range of 10–1000 nm) through the process of shear by pressing a sample through a die with two intersecting channels of equal cross sections (Fig. 3.65). The cross section remains unchanged by the repetitive pressing of the sample. A high total strain rate is achieved by the multiple pass pressing. Various routes are possible, depending on the orientations, which are designated as A, B, and C. (the sample is rotated 0, 90, and 180 degrees along its longitudinal axis). Routes designated as BA and Bc refer to a rotation of 90 degrees in the opposite and in the same sense between consecutive passes, respectively. The process yields a lamellar microstructure when their

Testing: Comparison of AM data with traditionally fabricated

115

Fig. 3.56 Several engineering stress-strain curves of specimens at the same loading conditions and the corresponding mean curve. (a) Compressive stress-strain curves of AM material; and (b) tensile stress-strain curves of wrought material at a higher strain rate. (Song, B., Nishida, E., Sanborn, B., Maguire, M., Adams, D., Carroll, J., Wise, J., Reedlunn, B., Bishop, J., Palmer, T., 2017. J. Dyn. Behav. Mater. 3, 412. With kind permission of Springer Nature.)

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Additive and traditionally manufactured components

Fig. 3.57 Comparison of microstructures of (a–c) wrought and (d–f) AM 304L stainless steel samples. Austenite grains are colored according to their orientation, while ferrite appears black in EBSD maps (a, b, d, e). The ferrite is identified with red arrows in the SEM images (c and f). Oxide inclusions in the AM material are highlighted with green circles. (Song, B., Nishida, E., Sanborn, B., Maguire, M., Adams, D., Carroll, J., Wise, J., Reedlunn, B., Bishop, J., Palmer, T., 2017. J. Dyn. Behav. Mater. 3, 412. With kind permission of Springer Nature.)

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Testing: Comparison of AM data with traditionally fabricated

Fig. 3.58 Schematic representation of orientations of the additively manufactured 304L stainless steel bars. (Song, B., Nishida, E., Sanborn, B., Maguire, M., Adams, D., Carroll, J., Wise, J., Reedlunn, B., Bishop, J., Palmer, T., 2017. J. Dynamic Behavior Mater. 3, 412. With kind permission of Springer Nature.)

Fig. 3.59 Dynamic compressive stress equilibrium. (Song, B., Nishida, E., Sanborn, B., Maguire, M., Adams, D., Carroll, J., Wise, J., Reedlunn, B., Bishop, J., Palmer, T., 2017. J. Dyn. Behav. Mater. 3, 412. With kind permission of Springer Nature.) Incident Bar Striker Bar

Double Pulse Shaper

Strain gage

Specimen

Transmission Bar

Momentum Bar

Strain gage

Fig. 3.60 A schematic representation of the Kolsky tension bar. (Song, B., Nishida, E., Sanborn, B., Maguire, M., Adams, D., Carroll, J., Wise, J., Reedlunn, B., Bishop, J., Palmer, T., 2017. J. Dyn. Behav. Mater. 3, 412. With kind permission of Springer Nature.)

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Additive and traditionally manufactured components

Fig. 3.61 Comparison of dynamic compressive stress-strain curves of wrought and AM/ Z-direction 304L stainless steel. (Song, B., Nishida, E., Sanborn, B., Maguire, M., Adams, D., Carroll, J., Wise, J., Reedlunn, B., Bishop, J., Palmer, T., 2017. J. Dyn. Behav. Mater. 3, 412. With kind permission of Springer Nature.)

Fig. 3.62 Comparison of dynamic compressive stress-strain curves of AM 304L stainless steel along annealed X and as-deposited Z-directions. (Song, B., Nishida, E., Sanborn, B., Maguire, M., Adams, D., Carroll, J., Wise, J., Reedlunn, B., Bishop, J., Palmer, T., 2017. J. Dyn. Behav. Mater. 3, 412. With kind permission of Springer Nature.)

Testing: Comparison of AM data with traditionally fabricated

119

Fig. 3.63 Comparison of dynamic compressive stress-strain curves of AM 304L stainless steel along the Z-direction but at different locations. (Song, B., Nishida, E., Sanborn, B., Maguire, M., Adams, D., Carroll, J., Wise, J., Reedlunn, B., Bishop, J., Palmer, T., 2017. J. Dyn. Behav. Mater. 3, 412. With kind permission of Springer Nature.)

orientation of the billet is not changed after each pass. Routes involving orthogonal deformation paths helps to achieve a microstructure with equiaxed grains. In Fig. 3.66 the orientation of compressive samples is indicated. The microstructures are seen in Fig. 3.67. In Fig. 3.67a the as-received initial microstructure is seen. The grains are approximately equiaxed and microstructure is homogeneous. The mean size is about 53 μm. A few annealing twins are observed. The optical microstructure of specimens undergoing 1 and 4 pass EPAC are seen in Fig. 3.67b and c, respectively. In Fig. 3.67d a microstructure of 304L SS is observed by TEM, which has undergone four passes resulting in grain size of 200–500 nm. The true stress-strain curves for orientations A, B, and C of the as-received 304L and for samples undergoing 1 and 4 passes ECAP are shown in Fig. 3.68. The strengths of the 304L SS increases after ECAP and are different for the orientations (loading directions) tested, but the work hardening rates are almost the same except in the four pass ECAP. There is thus an anisotropy induced by EPAC. This can be observed more

120

Table 3.23 Material processing information and testing conditions. Material

AM 304L; 2.0 kW Cross-hatched

Postprocessing

Strain rate (s21)

Specimen location within parent bar

X

Annealed

Z

None (asdeposited)

Longitudinal

None

Transverse

None

Compression 500 1500 3000 Tension 3000 Compression 500 500 1500 1500 3000 3000 Tension 3000 3000 Compression 500 1500 3000 Tension 3000 Compression 500 1500 3000 Tension 3000

Center Center Center Center Center Edge Center Edge Center Edge Center Edge N/A N/A N/A N/A N/A N/A N/A N/A

Song, B., Nishida, E., Sanborn, B., Maguire, M., Adams, D., Carroll, J., Wise J., Reedlunn, B., Bishop, J., Palmer, T., 2017. J. Dyn. Behav. Mater. 3, 412. With kind permission of Springer Nature.

Additive and traditionally manufactured components

Wrought 304L

Orientation of test sample load axis

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Testing: Comparison of AM data with traditionally fabricated

1000 298K; 0.001/s

True Stress (MPa)

800

600 1200 1100

400

1000 900 800

200

AM (As-built) Wrought AM (Recrystallized)

233K rate change 0.001/s 0.1/s

700 600 0.1

0.12

0.14

0.16

0.18

0.2

0 0

0.05

0.1

0.15

0.2

0.25

0.3

True Strain

Fig. 3.64 Compressive true stress-true strain of 316L SS in the AM-as-built condition, AM + recrystallization heat treatment, and annealed wrought 316L SS. (Gray III, G.T., Livescu, V., Rigg, P.A., Trujillo, C.P., Cady, C.M., Chen, S.R., Carpenter, J.S., Lienert, T.J., Fensin, S., 2015. EPJ Web of Conferences 94, 02006. Open access.)

Fig. 3.65 Optical microscopy of the 316L SS materials studied: (a) annealed wrought plate, (b) AM-(as-built), and (c) AM-following recrystallization heat treatment at 1060°C for 1 h (AM-Rx). (Gray III, G.T., Livescu, V., Rigg, P.A., Trujillo, C.P., Cady, C.M., Chen, S.R., Carpenter, J.S., Lienert, T.J., Fensin, S., 2015. EPJ Web of Conferences 94, 02006. Open access.)

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Additive and traditionally manufactured components

Fig. 3.66 Compressive sample orientations with respect to the ECAP billet geometry. The inclined lines indicate the shear flow plane from the one pass or four passes. The angles between extrusion direction and compression axis are equal to 0, 45, and 90 degrees, respectively. (Qu, S., Huang, C.X., Gao, Y.L., Yang, G., Wu, S.D., Zang, Q.S., Zhang, Z.F., 2008. Mater. Sci. Eng. A 475, 207. With kind permission of Elsevier.) (b)

(a)

100 mm

(c)

50 mm

(d)

50 mm

250 nm

Fig. 3.67 Microstructures of AISI 304L stainless steel observed on X plane: (a) asreceived; ECAPed for (b) one pass and (c) four passes observed by optical microscope; (d) four passes observed by TEM. (Qu, S., Huang, C.X., Gao, Y.L., Yang, G., Wu, S.D., Zang, Q.S., Zhang, Z.F., 2008. Mater. Sci. Eng. A 475, 207. With kind permission of Elsevier.)

123

400 350

(b) 1200 B B

300

True stress (MPa)

(a) True stress (MPa)

Testing: Comparison of AM data with traditionally fabricated

A C

250 200

A: 0° B: 45° C: 90°

150 100

C

1000 800

A A: 0°

600

B: 45° 400

C: 90°

200

50 0 0.00

0.01

0.02

0.03

0.04

0.05

0 0.00

0.06

0.02

0.04

True strain

0.06

0.08

0.10

0.12

True strain

(c) True stress (MPa)

1200

B

A

1000

C

800 600

A: 0°

400

B: 45° C: 90°

200 0 0.00

0.01

0.02

0.03

0.04

0.05

0.06

0.07

True strain

Yield stress (MPa)

(a) 1000

(b)

800 600 As-received One pass Four passes

400 200 0

0

15

30

45

60

Angle (degree)

75

90

Strain hardening rate ds/de (GPa)

Fig. 3.68 Compressive stress-strain curves of the AISI 304L stainless steel samples with different orientations: (a) as-received, ECAPed for (b) one pass and (c) four passes. (Qu, S., Huang, C.X., Gao, Y.L., Yang, G., Wu, S.D., Zang, Q.S., Zhang, Z.F., 2008. Mater. Sci. Eng. A 475, 207. With kind permission of Elsevier.)

7 As-received One pass Four passes

6

5

4

3 0

15

30

45

60

75

90

Angle (degree)

Fig. 3.69 Dependence of (a) yield stress and (b) strain hardening rate on the number of ECAP passes and the angle between the extrusion direction and the compression axis. (Qu, S., Huang, C.X., Gao, Y.L., Yang, G., Wu, S.D., Zang, Q.S., Zhang, Z.F., 2008. Mater. Sci. Eng. A 475, 207. With kind permission of Elsevier.)

pronouncedly in the yield strength variation with the extrusion direction relative to the loading axis. The strain hardening of the four pass EPAC is very anisotropic as seen in Fig. 3.69b, however, the yield stress anisotropy is more pronounced in the one pass EPAC. From the straight portion of

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Additive and traditionally manufactured components

Fig. 3.70 SEM micrograph of samples after compression: (a) as-received and ECAPed for four passes, (b) sample A, (c) sample B, and (d) sample C. (Qu, S., Huang, C.X., Gao, Y.L., Yang, G., Wu, S.D., Zang, Q.S., Zhang, Z.F., 2008. Mater. Sci. Eng. A 475, 207. With kind permission of Elsevier.)

these figures Δσ and Δε were determined and the linear strain hardening rate was evaluated using the relation dσ Δσ ¼ dε Δε

(3.39)

SEM micrographs of the 304L SS after the compression tests are shown in Fig. 3.70. It can be seen in Fig. 3.70a that after the compression test many slip bands are present in the grains of the as-received and EPCA treated specimen by four passes. The slip bands are parallel to each other. A streamline feature exists after four pass EPAS after the compression tests as seen in Fig. 3.70b–d. The streamline feature is in the lateral surface of the compressive samples and the streamline planes make different angles—indicated in the figure—with respect to the compressive direction for the samples A, B, and C, respectively. The streamline planes are almost consistent with the direction of the shear planes. The sample orientation determines the shear planes with respect to the compressive direction for the samples shown in Fig. 3.66. The approximate angles are 27, 18, and 63 degrees as illustrated in Fig. 3.71. Thus the shear stress τσ can be expressed as τs ¼ σ c sin θ cosθ

(3.40)

125

Testing: Comparison of AM data with traditionally fabricated

θ = 18°

θ = 63°

Loading direction

Shear bands

θ = 27°

(a)

(b)

(c)

Fig. 3.71 Illustration of the shear band distribution in the samples a, b, and c with an interaction angle of 27, 18, 63 degrees, respectively. Here, θ is the angle between the shear bands and the loading direction and φ is the angle between the compression axis and the extrusion direction. Sample a (φ ¼ 0 degrees), sample b (φ ¼ 45 degrees), and sample c (φ ¼ 90 degrees). (Qu, S., Huang, C.X., Gao, Y.L., Yang, G., Wu, S.D., Zang, Q.S., Zhang, Z.F., 2008. Mater. Sci. Eng. A 475, 207. With kind permission of Elsevier.)

where σ c is the compressive stress and θ is the angle between the shear plane and the loading direction (see Fig. 3.72). The angles for A, B, and C are indicated in Fig. 3.71, respectively. From Eq. (3.40) the compression stress, σ c is expressed as σc ¼

τs sin θ cos θ

(3.41)

Assuming that the critical shear strength, τs is nearly the same, the compressive stress should be a function of the angle θ between the shear plane and the loading direction. Thus substituting the angles θΑ ¼ 27 degrees, θB ¼ 18 degrees, and θC ¼ 63 degrees into Eq. (3.41) one gets three equations for σ c the compressive strength as τs ¼ 2:4τs sin 27o cos 27o τs σ Bc ¼ ¼ 3:4τs sin18o cos18o τs σC ¼ 2:4τs c ¼ o sin 63 cos 63o σ Ac ¼

(3.42)

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Additive and traditionally manufactured components

σc

τ

σθ θ

θ Shear plane

σc Fig. 3.72 Illustration of the normal and shear stresses on the shear plane of the AISI 304L stainless steel sample ECAPed for one pass under compressive test. (Qu, S., Huang, C.X., Gao, Y.L., Yang, G., Wu, S.D., Zang, Q.S., Zhang, Z.F., 2008. Mater. Sci. Eng. A 475, 207. With kind permission of Elsevier.)

Table 3.24 The yield stress and strain hardening rate of the as-received and the ECAPed samples with different angles with respect to the extrusion direction. Processing condition As-received

ECAP for one pass

ECAP for four passes

Angle Angle Angle Angle Angle Angle Angle Angle Angle 0° 45° 90° 0° 45° 90° 0° 45° 90°

Yield stress 150 (MPa) dσ/dε 3.50 (GPa)

199

176

652

877

850

768

881

832

3.70

3.50

5.20

5.29

5.12

3.90

5.80

4.30

Qu, S., Huang, C.X., Gao, Y.L., Yang, G., Wu, S.D., Zang, Q.S., Zhang Z.F., 2008. Mater. Sci. Eng. A 475, 207. With kind permission of Elsevier.

Table 3.24 summarizes the experimental data of sample exposed to the ECPA treatment. One can note that the ECAP technique causes the anisotropy observed in the mechanical properties as expressed by the compression tests. In particular, the anisotropy in the yield stress after one pass ECAP should

127

Testing: Comparison of AM data with traditionally fabricated

be noted (Fig. 3.69), although the anisotropy decreases significantly for the four pass ECAP.

3.2.6 Ceramics—Alumina Fabrication of ceramics is difficult due to its brittleness. 3D printing (also known as AM) is a method allowing three-dimensional production of ceramics through a layer-by-layer deposition process, and thus overcoming fabrication difficulties. A schematic illustration of the fabrication of three-dimensional objects by the ExOne M-Lab system is illustrated in Fig. 3.73. The system is composed of the two-bed system; the one is for the powder (here alumina), and the other is for the actual fabrication. The basic principle is to spread the powder with a roller while the bed is moving mechanically in the X-direction. Following the powder spreading uniformly the bed returns to its original position where a binder droplets are selectively deposited through an electric printhead for part fabrication. Following the binder deposition, the powder-bed moves below a heater for a certain time required for powder binding, which is a one layer fabrication (one full cycle), and the process repeats itself until the part fabrication is completed.

Binder Feed

Leveling Roller Powder Supply Feed Stage

Binder Jetting Head

Part Powder Bed Build Stage

Fig. 3.73 Schematic representation representing the ExOne components used for the fabrication process. (Gonzalez, J.A., Mireles, J., Lin, Y., Wicker, R.B., 2016. Ceram. Int. 42, 10559. With kind permission of Elsevier.)

128

Additive and traditionally manufactured components

100

100 320 Grit

80 60 26 mm

40

53 mm

20

106 mm

Relative ρ (%)

Relative ρ (%)

240 Grit 80 60

23 mm

40

45 mm

20

90 mm

0

0 2h 16h Sintering Profile

100

Relative ρ (%)

100

400 Grit

80 60

15 mm

40

30 mm

20

60 mm

Mixed Powder

80 60

15 mm

40

45 mm

20

106 mm

0

0

(C)

2h 16h Sintering Profile

(B)

Relative ρ (%)

(A)

2h 16h Sintering Profile

(D)

2h 16h Sintering Profile

Fig. 3.74 Relative density vs sintering profiles including the three-layer thickness used for each powder. (A) 240 grit size powder or 53 μm particle size, (B) 320 powder size or 45 μm particle size, (C) 400 grit size powder or 30 μm particle size, and (D) combination of all three powders. (Gonzalez, J.A., Mireles, J., Lin, Y., Wicker, R.B., 2016. Ceram. Int. 42, 10559. With kind permission of Elsevier.)

The parts were sintered for two dwell times, 2 and 16 h at temperatures up to 1600°C. The relative density profiles following sintering are illustrated in Fig. 3.74 for the two dwelling times. SEM images of the powder used for the AM process, the green body and those of the sintered parts are shown in Fig. 3.75. Sintering causes the powder particles to diffuse together decreasing the porosity and increasing the density. The 16-h sintered material is shown in Fig. 3.76 at higher magnification. Note the neck formation, which might indicate some elongation in the fully sintered alumina at this high-temperature sintering (1600°C). Compression tests were performed on the mixed powder parts sintered at both temperatures. The compression results are shown in Fig. 3.77. A nearly fully sintered (96%) alumina part could be achieved with an average compressive strength of 131.86 MPa when sintering dwell time was 16 h (Fig. 3.77B). Clearly, various build parameters, such as particle size, layer thickness or saturation determine the resulting properties—physical and mechanical alike—of the alumina part.

Testing: Comparison of AM data with traditionally fabricated

129

Fig. 3.75 SEM images for (A) as purchased powder (B) green body fabricated part (C) part sintered for 2 h and (D) part sintered for 16 h. All parts were mixed powders at 45-μm layer thickness. (Gonzalez, J.A., Mireles, J., Lin, Y., Wicker, R.B., 2016. Ceram. Int. 42, 10559. With kind permission of Elsevier.)

Fig. 3.76 SEM image of 16-h sintered part with a high magnified image of creation of neck between two powder particles. (Gonzalez, J.A., Mireles, J., Lin, Y., Wicker, R.B., 2016. Ceram. Int. 42, 10559. With kind permission of Elsevier.)

130

Additive and traditionally manufactured components

90

(A)

80 70

Stress (MPa)

60 50

Sample 1 Sample 2

40

Sample 3 Sample 4

30 20 10 0 0 0.25 0.5 0.75 1 1.25 1.5 1.75 2 2.25 2.5 2.75 3 3.25 Strain (mm/mm)

(B)

160

140

120

Stress (MPa)

100 Sample 1

80

Sample 2 Sample 3

60

Sample 4

40

20

0 0 0.25 0.5 0.75 1 1.25 1.5 1.75 2 2.25 2.5 2.75 3 3.25 3.5 Strain (mm/mm)

Fig. 3.77 Stress-strain curves of the alumina compression test. (A) Samples that were sintered for 2 h and (B) samples that were sintered at 16 h. (Gonzalez, J.A., Mireles, J., Lin, Y., Wicker, R.B., 2016. Ceram. Int. 42, 10559. With kind permission of Elsevier.)

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Testing: Comparison of AM data with traditionally fabricated

3.2.7 Conventionally fabricated alumina (Al2O3) Accuratus Ceramic Corporation indicates the compressive stress at room temperature of 99.5% alumina in the following table. Mechanical

Compressive strength

Units of measure

SI/ metric

(Imperial)

MPa (lb/in2  103)

2600

(377)

>2100

2500

315 375

675 330

Compressive strength (MPa) 2200–2600 Mechanical Compressive MPa >2100 strength Flexural strength MPa 370 Young’s modulus GPa 275

Goodfellow—in its many commercial publications—indicates for alumina typical data for the compressive strength value in the range of 2200– 2600 MPa. Clearly, the actual strength depends on many parameters the crucial of them is the porosity content. Porosities can develop into cracks. Therefore, a large body of experimental research activity was focused on the effect of pores on alumina. In Fig. 3.78 the compressive stress in porosive alumina is shown. The failure shape of the specimens after compression tests with different porosities are seen in Fig. 3.79. Note that the shapes are different. In samples with 18%–29% porosity (not shown in the figure) brittle fracture occurs of elastic specimens after elastic energy accumulation. Due to elastic energy release, the specimen fully breaks down after reaching the compression strength limit. With the increase of pores, the damage is more localized and the specimens do not fully break down after reaching the compression strength limit, therefore ability exists for further deformation. In Fig. 3.79a–c the specimens contain 40% porosity. Cracks are formed at 45 degrees to the applied load on the lateral surfaces of cylindrical specimens as indicated in Fig. 3.79a. Multiple cracks are formed during the deformation process when the porosity in the alumina ceramics is in the range 50%–70%. These cracks are parallel to the loading axis as illustrated in Fig. 3.79c located as mentioned earlier on the lateral surface of the cylindrical specimen. It could be mentioned—since it is included in Fig. 3.79—that in shear tests (Fig. 3.79e) in specimens with 40% porosity the main crack was at 45 degrees to the applied load and the crack appeared immediately after the elastic region of the stress-strain curve. Above 40% porosity the deformation was associated with many fine cracks in different directions until complete failure as seen in Fig. 3.79f. In each

132

Additive and traditionally manufactured components

400

12 10

300

σ, MPa

σ, MPa

8 -1

200

-1

6

-2 4

-2

100

2 0

0 0

0.004

0.008

0.012

0.016

0.02

0

0.02

ε

0.04

0.06

0.08

ε 14

4

12 10

σ, MPa

σ, MPa

3 -1 2

2 8 6 4

1

2 0

0 0

0.01

0.02

ε

0.03

0.04

0

0.02

0.04

0.06

0.08

ε

Fig. 3.78 Stress-strain diagrams obtained by tests in compression (curve 1) and shear (curve 2) for Al2O3 with porosity of: (a) 20% and (b) 50%. (c) Stress-strain diagram obtained by tests in compression for alumina ceramics with a pore space volume over 60%. (d) Stress-strain diagram obtained by tests in shear for alumina ceramics with a pore space volume over 60%. (Savchenko, N., Sevostyanova, I., Sablina, T., Go€mze, L., Kulkov, S., 2014. AIP Conference Proceedings 1623, 547. With kind permission of AIP Conference Proceedings.)

porous ceramic material (such as alumina), each pore is a potential source of microcracking and clearly with the increase in the pores the sources of microcracks also increases after compression (or shear). Thus, pores in the ceramic material effect the overall mechanical properties. 3.2.7.1 Effect of orientation and temperature In single-crystal alumina, known as sapphire the compressive and flexure strength depend on the orientation of the crystal and temperature. The decrease in the mechanical strength with temperature is usually rapid

Testing: Comparison of AM data with traditionally fabricated

133

Fig. 3.79 The failure shape of Аl2O3 specimens with porosities of (a and b) 40%, and (c and d) 55% after a test in compression. The failure shape of Аl2O3 specimens with porosities of: (e) 40%; (f) 69% after a test in shear. (Savchenko, N., Sevostyanova, I., Sablina, T., Go€mze, L., Kulkov, S., 2014. AIP Conference Proceedings 1623, 547. With kind permission of AIP Conference Proceedings.)

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Additive and traditionally manufactured components

2000

Compressive strength (a-axis)

Compressive strength (a-axis) (Hurley)

Strength (MPa)

1500

1000 Tensile strength (c-axis)

Compressive strength (c-axis)

Compressive strength (c-axis) (Hurley)

Diamonds: Scott & Orr data

500

Tensile strength (a-axis) 0 0

200

400

600

800

1000

1200

Temperature (∞C)

Fig. 3.80 Compressive and tensile strength of sapphire measured along the a- or c-axis of the crystal. (Schmid, F., Harris, D.C., 1998. J. Am. Ceram. Soc. 81, 885. With kind permission of John Wiley and Sons.)

compared to the ceramic materials. The c-axis compressive strength is lower than the tensile stress at temperatures 400°C. It decreases below 2% of the room temperature compressive strength at 800°C. Compression parallel to the c-axis induces twinning in the crystal, which leads to mechanical failure due to the intersection of twins on different rhombohedral planes. The effect of c-axis compression is shown in Fig. 3.80. Note that the compressive stress at the temperatures indicated in the a direction changes comparatively less than in the c direction where the drop in the compressive stress is very rapid. The tensile strength of the sapphire along the a- or c-axis almost does not change in the temperature region 20–800°C. Similarly, the compression stress changes little in the a direction up to 1300°C. The applied load against the crosshead displacement for compression (and tension) tests is shown in Fig. 3.81. As mentioned twinning is observed for the c-axis compression at temperatures 200°C, but was not observed in the a-axis compression although twinning was observed in a fractured specimen. Thus, the porosity of brittle materials can have a significant effect on their physical (mechanical, thermal, electrical) properties. The macroscopic behavior of ceramics can vary from completely brittle to quasi-plastic depending on the pore space volume.

135

Testing: Comparison of AM data with traditionally fabricated

c-Axis compression 16.2 14.1 2.54 0.44

0.32

a-Axis c-Axis tension tension 11.4 9.12

Load

Ultimate load (kN):

20°C 200°C

400°C 600°C 800°C Crosshead displacement

800°C

800°C

Fig. 3.81 Load versus crosshead displacement for compression and tension tests of the c-axis and a-axis sapphire. Note that the horizontal and vertical scales are different for each curve; the purpose of the illustration is to compare the shape of the curves. (Schmid, F., Harris, D.C., 1998. J. Am. Ceram. Soc. 81, 885. With kind permission of John Wiley and Sons.)

3.3 Indentation (hardness) As known indentation hardness measures the resistance to plastic deformation resulting from the application of an indenter at a constant load. The resistance to indentation is a consequence of strong intermolecular bonds in the solid material, but its behavior when under force is quite complex and requires skill for interpretation, despite the simplicity of the measurement. It is related to the flow stress by various empirical relations. The strength displayed by hardness is a measure of material’s elastic range or of both the elastic and plastic range together. The choice of the indenter depends on the material whether soft or hard. An important hard material is ceramics or ceramic—like substances. The most common hardness tests are performed by Rockwell, Vickers, and Brinnel. The choice of the indenter is determined by the tester.

3.3.1 Ti-6Al-4V Ti-6Al-4V often written as Ti6Al4V is an important engineering alloy with attractive properties and very good workability, which allows its fabrication

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Additive and traditionally manufactured components

into complex shapes. Due to its attractive properties, it has been extensively used in aircraft industries. One of the AM processes is by in situ alloying of wire fed Ti and Al on a pure Ti substrate to produce Ti6Al4V layer-by-layer fashion. This process is based on gas tungsten arc welding (GTAW) to produce the three-dimensional object. Al and Ti wires are delivered separately to the weld pool; thus in situ alloying is performed. The GTAW was used as a heating source to produce the melt pool into which the Ti and Al were delivered individually by two wire feeding units. Clearly, a new layer was deposited on a previous layer until the part (or component) was fully constructed as mentioned in the layer-by-layer fashion. After the process was finished as the end product was fabricated, the Ti substrate was removed. For the experimental details of the four groups of alloys, the original work could be consulted. In brief, four straight walls were produced by additive laser manufacturing (ALM) each of 100 mm length. The representative microstructure of the regions is illustrated in Fig. 3.82. Vickers values are shown in Fig. 3.83. In Fig. 3.83a the hardness variation is indicated with distance from the fusion line, while in Fig. 3.83b the average Vickers hardness with wire feed ratio are shown. In Fig. 3.83a the effect of the feed ratio is indicated. The cooling rate is a significant factor that controls the proportion of the α2-phase which has a hardness of twice of the γ. These phases comprise the microstructure illustrated in Fig. 3.82. It has been indicated earlier that various empirical relations were developed for evaluating stress (e.g., tensile) from hardness measurements. In many alloys and metals a strong correlation exists between tensile strength and hardness; in Ti6Al4V limited AM data are available, but recently relation between hardness and tensile stress in Ti6Al4V appeared and reproduced in Fig. 3.84. The Ti6Al4V alloy was obtained by powder-bed AM and the microstructure is shown in Fig. 3.85—a fully acicular martensite of α0 and growth bands. There is an influence of heat treatment time and temperature on the microstructure. The α0 is converted to a mixture of lamellar α and β on heating below the β transus (see, e.g., the Ti-Al phase diagram in Fig. 6 of Tong et al. (2017)) but features of the original microstructure are retained. On heating above the β transus large grain growth occurs, but the α0 still transforms to the lamellar α and β. Heating at 850°C for 2 h, for example, and air cooling induces α-Widmanst€aten (basket weave structure) formation as seen in Fig. 3.85b, which is different than that found in the wrought alloy seen in Fig. 3.86a where isolated lamellae of α and β are observed between almost equiaxed grains. The microstructure of the AM Ti6Al4V also consists of martensitic α0 -phase as seen in Fig. 3.86b and c. After subjecting the alloy to HIP at 900°C for 2 h at 102 MPa as shown in Fig. 3.86d, the structure is

Testing: Comparison of AM data with traditionally fabricated

137

Fig. 3.82 Representative microstructure in the wall components produced by GTAW ALM: (a) top region, the white areas show the interdendritic γ-phase; (b) middle region, the white areas show the interdendritic γ-phase; (c) middle region, lamellar colonies consist of α2 and γ lamellae; and (d) bottom region, coarse α2 grains with fine γ laths precipitating at grain boundaries. (Ma, Y., Cuiuri, D., Hoye, N., Li, H., Pan, Z., 2014. J. Mater. Res. 29, 2066. With kind permission of Cambridge University Press.)

similar but coarser than the heat-treated alloy as seen in Fig. 3.85b. LBM and EBM are laser and electron beam melting acronyms. One is puzzled if the material properties of parts manufactured by AM are equivalent to those of the wrought one since there are distinct differences in the microstructures. Tong et al. claim that the tensile and other mechanical properties such as yield or ultimate tensile strength are comparable or even sometimes better than the wrought alloy. 3.3.1.1 Conventionally produced Ti-6Al-4V The microstructure and hardness of Ti6Al4V are presented in this section and they are compared with similar properties of commercially pure Ti. Such a comparison indicates the preferential use of Ti6Al4V for aero and automotive industries. For start, Table 3.25 summarizes and compares the

138

Additive and traditionally manufactured components

550 0.80 0.95 1.12

r-s ea N e

t ra

st ub

500

1.30

Vickers Hardness

450 400 350 300 250 200

0

(a)

2

4

6

8

10

12

14

16

18

20

22

24

26

Distance from the fusion line (mm)

Average Vickers Hardness (HV0.2)

400

350

300

250

0.80 0.85 0.90 0.95 1.00 1.05 1.10 1.15 1.20 1.25 1.30 1.35

(b)

Wire feed rate ratio

Fig. 3.83 Microhardness values of wall components produced by GTAW ALM under different wire feed conditions. (a) Microhardness profiles as a function of location on the cross sections. (b) Mean microhardness values of the main part except for the nearsubstrate zone. Error bars show one standard deviation. (Ma, Y., Cuiuri, D., Hoye, N., Li, H., Pan, Z., 2014. J. Mater. Res. 29, 2066. With kind permission of Cambridge University Press.)

microhardness values of Ti6Al4V and Ti. The average of five trials is included in the last column of the table. Note that the Vickers microhardness is much higher than that of pure Ti. The Rockwell C data listed are only for the Ti6Al4V. The macrohardness of these alloys is compared in Table 3.26.

Testing: Comparison of AM data with traditionally fabricated

139

Fig. 3.84 Hardness and tensile strength. (Tong, J., Bowen, C.R., Persson, J., Plummer, A., 2017. Mater. Sci. Tech. 33, 138. With permission of Taylor and Francis.)

Fig. 3.85 (a) Columnar martensitic microstructure in as-built Ti-6Al-4V, (b) 2 h at 850°C air cooling. The α-phase is light and β is dark. (Tong, J., Bowen, C.R., Persson, J., Plummer, A., 2017. Mater. Sci. Tech. 33, 138. With permission of Taylor and Francis.)

140

Additive and traditionally manufactured components

Fig. 3.86 Microstructure of Ti-6Al-4V (a) wrought, (b) horizontal powder-bed fusion, (c) vertical powder-bed fusion, and (d) horizontal after HIP (900°C/102 MPa/2 h). (Tong, J., Bowen, C.R., Persson, J., Plummer, A., 2017. Mater. Sci. Tech. 33, 138. With permission of Taylor and Francis.)

Table 3.25 A compilation of microhardness test data made on the two materials Ti-6Al-4V alloy and commercially pure titanium (Grade 2). Alloy

Trial 1

Trial 2

Trial 3

Trial 4

Trial 5

Average

Sample 2ATi-6Al-4V Vickers hardness Rc Sample 2B,Ti-6Al-4V Vickers hardness Rc Sample 1A, CP Grade 2 Vickers hardness Rc Sample 1B, CP Grade 2 Vickers hardness Rc

340.41 325.20 332.72 334.61 340.46 334.68 25 27 – 30 25 26.75 330.88 342.44 332.74 327.11 338.51 334.36 28 25 25 27 28 26.6 213.59 195.36 193.65 193.67 200.58 199.37 – – – – – – 175.59 178.58 167.21 169.26 193.67 176.82 –

–

–

–

–

–

Grade 2 refers to the commercial Ti. Poondla, N., Srivatsan, T.S., Patnaik, A., Petraroli, M., 2009. J. Alloys Compounds 486, 162. With kind permission of Elsevier.

141

Testing: Comparison of AM data with traditionally fabricated

Table 3.26 A compilation of macrohardness test data made on the two materials Ti-6Al-4V alloy and commercially pure titanium (Grade 2). Alloy

Trial 1

Trial 2

Trial 3

Trial 4

Trial 5

Average

Sample 2A, Ti-6Al-4V Vickers hardness Rc Sample 2B, Ti-6Al-4V Vickers hardness Rc Sample 1A, CP Grade 2 Vickers hardness Rc Sample 1B, CP Grade 2 Vickers hardness Rc

286

286

302

279

286

287.8

28 286

28 294

30 310

27 279

28 286

28.2 291

28 260

29 260

31 272

27 254

28 260

28.6 261.2

24 260

24 266

26 279

23 254

24 260

24.2 263.8

24

25

27

23

24

24.6

Poondla, N., Srivatsan, T.S., Patnaik, A., Petraroli, M., 2009. J. Alloys Compounds 486, 162. With kind permission of Elsevier.

Note that almost no difference or very small difference is recorded for the two alloys, that of Ti6Al4V and Ti. Further, it is likely that the difference between the macrohardness and the microhardness values is associated with the population, distribution, and amount of the relative phases. In the Ti6Al4V alloy, two phases are present primary α and β in the lamellar α and β region. Apparently, the indenter in the macrohardness test feels a large surface area containing a large portion of the α +β lamellar region containing the softer β-phase compared to the microindenter (see, e.g., Figs. 3.87b and 3.89b). This is supported by the fact that the phase of Ti is relatively stronger in strength than the β-phase. During the microindentation, the possibility exists to hit larger amounts of the stronger (harder) α provided the indentation was done randomly on the polished surface (and not selectively). Generally, β-phase titanium is the more ductile phase (softer) and α-phase is stronger and less ductile (harder). The above speculation for the indicated difference is one possibility among others to explain the values obtained by the microhardness and macrohardness assuming the same treatment of the specimens tested. The microstructure affects significantly the mechanical properties including the magnitude of the hardness. Therefore, before comparing the hardness values measured, it is of interest to look at the microstructures. In Fig. 3.87, a comparison of microstructures is made between Ti and Ti6Al4V at the same magnification. Also, the microstructures of Ti and Ti6Al4V are shown for comparison by optical micrographs at various magnification in Figs. 3.88 and 3.89. The hardness values for Ti and Ti6Al4V

(a)

(b)

10 mm

10 mm

Fig. 3.87 Optical micrographs comparing the intrinsic microstructural features of the two materials, commercially pure (Grade 2) and Ti-6Al-4V at equivalent magnification of 1000. (Poondla, N., Srivatsan, T.S., Patnaik, A., Petraroli, M., 2009. J. Alloys Compounds 486, 162. With kind permission of Elsevier.)

Fig. 3.88 Optical micrographs showing the key microconstituents in the commercially pure (Grade 2) titanium at three different magnifications. (Poondla, N., Srivatsan, T.S., Patnaik, A., Petraroli, M., 2009. J. Alloys Compounds 486, 162. With kind permission of Elsevier.)

Testing: Comparison of AM data with traditionally fabricated

143

Fig. 3.89 Optical micrographs showing the key microconstituents, their morphology and distribution, in the Ti-6Al-4V alloy at three different magnifications. (Poondla, N., Srivatsan, T.S., Patnaik, A., Petraroli, M., 2009. J. Alloys Compounds 486, 162. With kind permission of Elsevier.)

in the form of bar graph is illustrated in Fig. 3.90. The Vickers hardness vs the number of measurements of both the microhardness and the macrohardness for Ti and Ti6Al4V are compared in Figs. 3.91 and 3.92.

3.3.2 Aluminum alloy (Al6061) 6061 Wires that were produced by friction extrusion served as the feedstock in a wire arc additive process (WAAM). The deposits were evaluated by microstructure, defect content, and strength. Currently, we are considering the hardness which was evaluated before and post build aging treatments.

144

Additive and traditionally manufactured components

Fig. 3.90 Bar graph depicting the average microhardness and macrohardness values of the commercially pure (Grade 2) and Ti-6Al-4V alloy. (Poondla, N., Srivatsan, T.S., Patnaik, A., Petraroli, M., 2009. J. Alloys Compounds 486, 162. With kind permission of Elsevier.)

This additive manufacturing method combines an electric arc heating source with a wire feeding system. It is a high deposition rate process. The Vickers hardness was measured in the transverse cross section along a central vertical line. Testing of the Al alloy was done in the as-build condition and after a T-6 aging treatment. The feedstock (friction extruded wires with different diameters) was used for the Gas Tungsten Arc (GTA)-based WAAM process. The experimental system is shown in Fig. 3.93. It consists of a power source, a robot, a TIG torch (3.2 mm electrode of W-2% Ce), and pure Ar shielding gas. The Vickers hardness of parts 4 and 5 fabricated by the WAAM process are shown in two conditions in the as-made and after 6061-T6 heat treatment in Fig. 3.94. The trend of the hardness variation for both cases is about the same, but the T6 heat treatment improved somewhat the hardness of the 6061 alloy. Parts 4 and 5 refer to the height of the walls made with wires number 4 and 5. The heights of these walls are 10 and 16 mm. Specimens from part 5 were machined for tensile specimen measurements. The hardness of part 5 and the UTS are compared as indicated in Fig. 3.95. Detailed discussion and results of the TIG-based WAAM process can be found in the work of Li et al. (2015b). As can be seen the hardness and the UTS show the same trend with distance, i.e., with the position in the build. In order to eliminate build defects, the wire quality should be good

145

Testing: Comparison of AM data with traditionally fabricated

Vickers Hardness-Hv(Kg/mm2)

(a) 250

200

150 0

1

2 3 4 Measurement Number

5

6

Vickers Hardness-Hv(Kg/mm2)

(b) 300

250

200

150 0

1

2

3

4

5

6

Measurement Number

Fig. 3.91 A profile showing the hardness values across the length of annealed commercially pure (Grade 2) titanium: (a) microhardness and (b) macrohardness. (Poondla, N., Srivatsan, T.S., Patnaik, A., Petraroli, M., 2009. J. Alloys Compounds 486, 162. With kind permission of Elsevier.)

for the automatic wire feeding in this process. The maximum hardness and tensile stress obtained by the WAAM process is as good as those of the AA6061 base metal (conventional fabrication). As mentioned, the surface condition of the wire should be good in order to eliminate voids (pores) formation in the WAAM process. Often in AM such as laser powder-bed fusion (LPBD) also known as selective laser melting (SLM) voids and cracks develop due to volumetric solidification shrinkage and thermal contraction. Such defects were illustrated in Fig. 3.8. High-temperature heating in laser powder-bed fusion

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Additive and traditionally manufactured components

Vickers Hardness-Hv(Kg/mm2)

(a) 400

350

300

250 0

Vickers Hardness-Hv(Kg/mm2)

(b)

1

2 3 4 Measurement Number

5

6

350

300

250 0

1

2 3 4 Measurement Number

5

6

Fig. 3.92 A profile showing the hardness values across the length of fully annealed Ti-6Al-4V alloy: (a) microhardness and (b) macrohardness. (Poondla, N., Srivatsan, T.S., Patnaik, A., Petraroli, M., 2009. J. Alloys Compounds 486, 162. With kind permission of Elsevier.)

additive manufacturing eliminates these defects as was illustrated in Fig. 3.9. Heating the power bed at 500°C effectively removes the defects. The microhardness values of LPBF produced AA 6061 are shown in Table 3.27. The results can be summarized as follows: Alloys with large solidification range, which are prone to appreciable shrinkage—such as AA 6061 which are inherently susceptible for defect formation in the LPBD AM technique, can be fabricated by this technique defect-free when hightemperature heating is added to the process. The mechanical properties including hardness of the heat-treated parts fabricated by LPBD values are comparable to annealed and heat-treated wrought alloys.

147

Testing: Comparison of AM data with traditionally fabricated

Robot Shielding gas

AC-TIG machine TIG torch

Substrate

Fig. 3.93 AC-Tig experimental system for wire and arc additive manufacturing. (Li, X., Reynolds, A.P., Baoqiang, C., Jialuo, D., Williams, S., 2015. TMS Annual Spring Meeting Supplemental Proceeding, TMS 2015. With kind permission of Springer Nature, p. 445.)

3.3.3 Conventionally fabricated Al 6061 Pure aluminum is soft and, therefore, may not be ideal for industrial applications. The Al 6061 is among the popular alloys in engineering applications, because of the properties can be induced by various heat treatments. Below, the effect of annealing on hardness and microstructure of cryorolled (CR) 6061 is presented. Cryorolling is a severe plastic deformation process used to obtain ultrafine-grained alloy with higher strength and hardness than in conventional cold rolling. To improve its formability warm forming after cryorolling may be applied. Isothermal heat treatment in 100–500°C temperature range showed the release of the lattice strain with the dissolution of precipitates and grain growth. Suppression of dynamic recovery following low-temperature aging has received great interest because of the high strength and very good ductility obtained after the treatment. In Fig. 3.96 optical microstructure and electron beam backscattered diffraction (EBSD) micrographs of as-received Al 6061 alloy are shown after solution treatment at 520°C for 2 h and after cryorolling up to 92% reduction. The solution treated (ST) alloy exhibits equiaxed grains an average of 80 mm in which dendritic segregation is observed (Fig. 3.96a). Low angle

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Additive and traditionally manufactured components

WAAM Part No.4 110 100 90 HV

80 70

As m

60

T6

50 40 WAAM Part No.5 110 100 90 HV

80 70

As made

60

T6

50 40 0

5

10

15

20

Distance from top of substrate [mm]

Fig. 3.94 Vickers hardness on vertical central line of transverse cross sections of part No. 4 and No. 5. (After Li, X., Reynolds, A.P., Baoqiang, C., Jialuo, D., Williams, S., 2015. TMS Annual Spring Meeting Supplemental Proceeding, TMS 2015. With kind permission of Springer Nature, p. 445.)

grain boundaries of 1.5–15 degrees and high angle grain boundaries >15 degrees are observed in Fig. 3.96C. The fraction of low angle grain boundaries observed in CR 92% reduction samples is 0.71. TEM micrographs of Al 6061 annealed at the temperatures indicated in the figure can be seen in Fig. 3.97. The Vickers hardness variation with temperature is shown in Fig. 3.98. The numbers on the graph indicate the hardness values at the specific temperatures. After the peak hardness at 150°C, a drastic softening occurs and in the range of 300–350°C the hardness drops to about 34–56 Vickers hardness. The alloy is a precipitation-hardening alloy. Thus, the improvement in hardness (and strength) at low-temperature annealing of 150°C is due to precipitation hardening. The precipitate is Mg2Si,

149

Testing: Comparison of AM data with traditionally fabricated

200 180 160 140 HV

120 100 80

As-made

60

T6

40 20 0 0

5

10

15

20

25

30

35

Distance from top of substrate [mm] 400 350

UTS [MPa]

300 250 200 As-made

150

T6

100 50 0 0

5

10

15

20

25

30

35

Distance from top of substrate [mm]

Fig. 3.95 Ultimate tensile stress and Vickers hardness of WAAM sample No. 5. (After Li, X., Reynolds, A.P., Baoqiang, C., Jialuo, D., Williams, S., 2015. TMS Annual Spring Meeting Supplemental Proceeding, TMS 2015. With kind permission of Springer Nature, p. 445.)

Table 3.27 Microhardness values of LPBF fabricated AA6061 components under different conditions. Condition

Microhardness

As built without powder-bed heating As built with powder-bed heating at 500°C T6 Heat-treated sample after fabrication with powder-bed heating at 500°C

90  6 HV 54  2.5 HV 119  6 HV

Uddin, S.Z., Murr, L.E., Terrazas, C.A., Morton, P., Roberson, D.A., Wicker, R.B., 2018. Addit. Manuf. 22, 405. With kind permission of Elsevier.

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Additive and traditionally manufactured components

Fig. 3.96 Optical and EBSD images of CR Al 6061 alloy. (A) optical image of ST at 520°C for 2 h; (B) optical image of 92% CR sample; (C) EBSD image of 92% CR sample; (D) gain boundary misorientation distribution of CR 92% sample. (Rao, P.N., Singh, D., Jayaganthan, R., 2013. Mater. Sci. Tech. 29, 76. With kind permission of Taylor and Francis.)

but the second precipitate of AlFeSi was also observed. The drop in hardness (and strength) on annealing at 200–300°C is a consequence of nucleation of dislocation free grains by recrystallization and their volume fraction increase. Additional hardness data of deformed 6061 alloy by accumulative role bonding (ARB) are shown in Fig. 3.99. The figure shows Vickers hardness variation with aging time after solution treatment and those of ARB processed specimens. Mg-Si precipitate clusters are considered to be the reason for the hardening at 100°C. The TEM micrograph of the ARB processed specimen is depicted in Fig. 3.100, while the specimens ARB processed and then aged is shown in the TEM illustration in Fig. 3.101. Fig. 3.101 shows two cases, (a) aged at 170°C for 3  104 s and (b) aged at 100°C for 2  105 s following the ARB process. Lamellar boundary structures are shown in the ARB proceed

Testing: Comparison of AM data with traditionally fabricated

151

Fig. 3.97 4 Images (TEM) of CR Al 6061 alloy annealed at different temperatures for 1 h. (A) 0°C; (B) 150°C; (C) 200°C; (D) 250°C (Rao, P.N., Singh, D., Jayaganthan, R., 2013. Mater. Sci. Tech. 29, 76. With kind permission of Taylor and Francis.)

(Fig. 3.100) and aged specimens (Fig. 3.101). Compared to pure Al, the ultrafine-grained microstructure of ARB fabricated Al6061 remained stable even during aging at 170°C and no growth has occurred. It can be seen in Fig. 3.99b that aging at low temperatures could improve the mechanical properties (illustrated by data of hardness) of Al alloys as a consequence of grain refinement and precipitation hardening. Thus, ARB deformed (room temperature) and solution treated specimens aged subsequently, age hardens with time at relatively low temperatures as illustrated in Fig. 3.99 with simultaneous changes in the microstructure.

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Additive and traditionally manufactured components

Fig. 3.98 Variation of hardness of CR Al 6061 alloy with annealing temperature. (Rao, P.N., Singh, D., Jayaganthan, R., 2013. Mater. Sci. Tech. 29, 76. With kind permission of Taylor and Francis.)

(a)

(b) 180

120

Vickers hardness, HV

Vickers hardness, HV

140

100 80 60 40

ST

20

ARB 6 cycles

0 As-ST or As-ARB

2

10

3

10

10

4

5

10

Aging time, t / s

10

6

160 140 120 100 80 60 40

ST

20

ARB 6 cycles

0 As-ST or As-ARB

10

2

10

3

10

4

105

106

Aging time, t / s

Fig. 3.99 Changes in Vickers hardness during aging of the ST specimen and the specimens ARB-processed by sixcycles as a function of aging time at (a) 170°C and (b) 100°C. (Terada, D., Kaneda, Y., Horita, Z., Matsuda, K., Hirosawa, S., Tsuji, N., 2014. IOP Conf. Series: Mater. Sci. Eng. 63, 012088. Open access.)

3.3.4 Stainless steel 304L Despite the extensive research activity and industrial use of 3D 304L and similar stainless steels (e.g., SS316; differing mainly in the 2% Mo addition), very limited direct information on the hardness of 3D 304L is reported in the literature. Most of the mechanical properties describing the strength of the 304L austenitic stainless steel pivot around tensile (and UTS) and yield strengths evaluation. As known, the austenitic stainless steel 304L is also known in Europe as 1.4307 stainless steel and is also sometimes termed as

153

Testing: Comparison of AM data with traditionally fabricated

RD ND 500nm

Fig. 3.100 TEM microstructure of the specimen ARB-processed by six cycles. (Terada, D., Kaneda, Y., Horita, Z., Matsuda, K., Hirosawa, S., Tsuji, N., 2014. IOP Conf. Series: Mater. Sci. Eng. 63, 012088. Open access.)

(a)

(b)

RD

RD

ND 500nm

ND 500nm

Fig. 3.101 TEM micrographs of the specimens ARB-processed by six cycles and then aged at (a) 170°C for 3  104 s and (b) 100°C for 2  105 s. (Terada, D., Kaneda, Y., Horita, Z., Matsuda, K., Hirosawa, S., Tsuji, N., 2014. IOP Conf. Series: Mater. Sci. Eng. 63, 012088. Open access.)

18/8 which indicates the Cr/Ni content. Each of these type steels is ironbased and alloyed with at least 10.5% chromium, which is what gives the metal its corrosion resistance. Austenitic stainless steel (ASS) 304L is being extensively used in the field of defense and nuclear science due to its excellent corrosion resistance in sea-water environment. This family of steels, the so-called 300 series are austenitic in nature because of the presence of Cr (16%–22%) and Ni (8%–14%). This range of Cr and Ni is found in similar grade steels such as SS316 which contains in addition 2% Mo for additional corrosion resistance. Indirect information on the hardness in 304L stainless steel can be obtained from a functionally graded Ti6Al4V to 304L component with a vanadium interlayer fabricated by laser metal deposition additive

154

Additive and traditionally manufactured components

Fig. 3.102 Schematic representation of RPM laser metal deposition system. (Reichardt, A., Dillon, R.P., Borgonia, J.P., Shapiro, A.A., McEnerney, B.W., Momose, T., Hosemann, P., 2016. Mater. Design 104, 404. With kind permission of Materials and Design.)

manufacturing. In the process, a melt pool is formed by rastering the laser beam across the sample surface, and powder is injected into the melt pool to deposit each layer. In the case considered, namely functional composition gradient, dissimilar material is deposited by layering an alloy directly on a dissimilar material. The gradient components were additively manufactured using a four hopper RPM (RPM and associates, In.) laser deposition system shown schematically in Fig. 3.102. The Vickers hardness measurement along the length of the component of interest, namely SS 304L is shown in Fig. 3.103. The hardness through the pure component of 304L is quite constant at a level of 185 Hv. The Vickers hardness indicated in Fig. 3.103 was confirmed in an earlier work evaluating the mechanical properties in 304L stainless steel also fabricated by additive manufacturing. This can be seen in Table 3.28. Note, however, that the Vickers hardness is layer thickness-dependent, decreasing with layer thickness increase. The samples were built using selective laser melting (SLM). Recall that it is an additive manufacturing technique that uses a laser to selectively melt powder particles layer-by-layer to build the product of interest.

155

Testing: Comparison of AM data with traditionally fabricated

Fig. 3.103 Vickers hardness results along the composition gradient of component B (i.e., 304L). (Reichardt, A., Dillon, R.P., Borgonia, J.P., Shapiro, A.A., McEnerney, B.W., Momose, T., Hosemann, P., 2016. Mater. Design 104, 404. With kind permission of Materials and Design.) Table 3.28 Mechanical properties for different layer thicknesses and scanning speeds. Layer thickness (μm)

30

Laser scanning speed (mm/s) Breaking elongation (%) Yield strength (MPa) Ultimate strength (MPa) Vickers hardness (10 kg) HV Micro Vickers (100 g) μ ΗV Micro Vickers (200 g) μ ΗV

70 25.9 182 393 192 211 217

50

90 22.1 156 389 196 205 209

70 10.9 146 383 156 203 212

70

90 11.2 142 377 164 198 207

70 8.6 139 334 159 192 191

90 4.2 132 316 144 189 193

Notes: Baseline for annealed, wrought 304L: yield strength—170 MPa, ultimate strength—480 MPa, Vickers hardness—185 max; breaking elongation—40% (ASTM Metals Handbook). Abd-Elghany, K., Bourell, D.L., 2012. Rapid Prototyp. J. 18, 420. With kind permission of Emerald Publishing Limited.

3.3.4.1 Conventionally produced 304L stainless steel The concept of ECAP has been mentioned earlier as an equal channel angular method of fabrication. In Fig. 3.104, the hardness in 304L stainless steel is shown. One can note that the hardness tends to increase with the numbers of the EPAC passes.

156

Additive and traditionally manufactured components

Vickers hardness (GPa)

4

3

2

1

0

0

1

2

3

4

Number of ECAP passes

Fig. 3.104 Dependence of Vickers hardness on the number of ECAP passes of the AISI 304L stainless steel. (Qu, S., Huang, C.X., Gao, Y.L., Yang, G., Wu, S.D., Zang, Q.S., Zhang, Z.F., 2008. Mater. Sci. Eng. A 475, 207. With kind permission of Elsevier.)

The mechanical properties including the hardness is seen in Table 3.29. The hardness is expressed in terms of GPa values. The ratio of the Vickers hardness to strength and yield stress is shown in Fig. 3.105. The well-known relation (Chen, 1980) between Vickers hardness and strength of σ f  HV/3 (σ f is the fracture strength) is not obtained for the as-received specimen. The ratios of HV/σ y and HV/σ f are equal to 7.9 and 1.7 for the yield stress and fracture strength respectively. On the other hand, the ECAP fabricated 1 and 4 pass specimens have almost the exact HV/σ f and HV/σ y values of 3. Table 3.29 Summary of the microstructure observation and results of the mechanical experiments conducted at room temperature. Compression (specimen A)

Tension Processing condition

UTS σ y (MPa) (MPa)

dσ/dε ε (%) (GPa)

dσ/dε σ y (MPa) (GPa)

Hardness (GPa)

As-received ECAP, one pass ECAP, four passes

168 710

556 790

61 34

1.473 1.020

150 652

3.5 5.2

1.33 2.21

1121

1136

12

0.175

768

3.9

3.02

Note: σ y, yield strength; UTS, ultimate tensile strength; ε, tensile fracture strain; dσ/dε, strain hardening rate. Qu, S., Huang, C.X., Gao, Y.L., Yang, G., Wu, S.D., Zang, Q.S., Zhang, Z.F., 2008. Mater. Sci. Eng. A 475, 207. With kind permission of Elsevier.

157

Ratio of hardness to strength

Testing: Comparison of AM data with traditionally fabricated

8

Hv / s y Hv / s f

6

4

2

0 0

1

4

The number of ECAP passes

Fig. 3.105 Relationship between HV and σ y(σ f) for as-received and ECAPed AISI 304L stainless steel samples. (Qu, S., Huang, C.X., Gao, Y.L., Yang, G., Wu, S.D., Zang, Q.S., Zhang, Z.F., 2008. Mater. Sci. Eng. A 475, 207. With kind permission of Elsevier.)

The hardness ratio of the as-received 304L can thus be written as σ y < HV/ 3 < σ f as seen in Fig. 3.105. It can thus be assumed that the hardness can be expressed as a function of the yield strength, σ y and the fracture strength σ f as HV ¼ xσ y + ð1  xÞσ f 3

(3.43)

In Eq. (3.43) σ y and σ f represent the contributions of the yield strength and the fracture strength to hardness, respectively. For the present case, i.e., as received 304L stainless steel x ¼ 0.34 calculated from Eq. (3.43) and Table 3.29. Many steel producers advertise their products providing data on the properties. Atlas Steels Grade 304 lists the mechanical properties of 304L as shown in the reproduced table in Table 3.30. The table also contains the hardness values for the 304L steel as the Brinnel hardness of HB ¼ 201. Except at high hardness values, the Vickers hardness and the Brinell hardness are just about identical as long as the Brinell impressions are of normal depth. Thus, 201 HB of the table is about the same value as also in HV.

3.3.5 Alumina Complex ceramic structures of complex and intricate shapes are difficult to fabricate. Al2O3 presented as an example manufactured by conventional

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Additive and traditionally manufactured components

Table 3.30 Mechanical properties of grade 304 steels. Mechanical property specification (single values are minima except as noted) Hardness

Tensile strength (MPa) Grade min

Yield strength 0.2% proof (MPa) min

Elongation (% in 50 mm) min

Rockwell B Brinell (HR B) (HB) max max

304 515 304L 485 304H 515

205 170 205

40 40 40

92 92 92

201 201 201

304H also has a requirement for a grain size of ASTM No 7 or coarser. AZO Materials, Stainless Steel—Grade 304.

technique requires large tool expenditure and long processing time. The new technique of relatively recent years, namely AM, has the definite advantage of fabricating complex parts with excellent dimensional stability of fine geometrical features some of which are difficult or even impossible to be produced by traditional methods. High purity alumina is used in bioceramics for hip prosthesis or tooth implants, etc. to mention some, which maintain their high hardness in use. AM based on stereolithography was used to prepare the green Al2O3. Sintering for densification of the slurry containing alumina powder and a mixture of organic binders (methylenbissacrylam is, acrylamide and glycerol) was performed after ball-milling for 6 h. SEM micrographs of the green aggregate and the thermal debinded structures are presented in Fig. 3.106. Debinding and sintering usually are associated with shrinkage. The changes in the average shrinkage at each stage of the debinding and sintering is shown in Fig. 3.107. The average shrinkage of samples in the length, width, and height are 2.65%, 2.61%, and 2.14%, respectively, as seen in Fig. 3.107A. The overall shrinkage is very small in all directions, indicating that the debinding had no apparent effect on the dimensions of the 3D oriented structures. The sintering shrinkage as a function of temperature is seen in Fig. 3.107B and its highest value is at 1650°C. The reason for the increase in shrinkage with temperature is associated with the volume decrease of pores with increasing temperature (atom rearrangement with temperature increase). The microstructural changes of the Al2O3 with increasing temperatures are illustrated in Fig. 3.108. As seen in the figure small rearrangement of grains occurred at 1200°C as a result of some necking among particles and no grain growth was observed in the temperature range 1200–1400°C but some pores between the alumina grains could be observed. At 1400°C or above

Testing: Comparison of AM data with traditionally fabricated

159

Fig. 3.106 SEM images of microstructure of green body (A) and the thermal-debinded body (B and C). (An, D, Li, H., Xie, Z., Zhu, T., Luo, X., Shen, Z., Ma, J., 1917. Int. J. Appl. Ceram. Tech. 14, 836. With kind permission of John Wiley and Sons.)

grain coarsening occurs due to the increase in diffusion rate inducing densification. With temperature increase at about 1600°C the grains were sintered together and the pores decreased as a result of solid-state diffusion. The neighboring layers were connected and the sintered samples had a homogeneous microstructure. The microhardness change (by Vickers hardness) with temperature as the structure densificates can be seen in Fig. 3.109. Table 3.31 compares the Al2O3 hardness obtained by 3D printing with other traditional methods. In Fig. 3.109 hardness with sintering temperature is shown. The microhardness increases with temperature up to 1600–1650°C resulting in 17.7–17.9 GPa, respectively, and is very much affected by porosity and grain size. The samples were not fully dense until 1600°C was achieved. Fig. 3.108 indicates the better AM hardness (and strength) of Al2O3 than the traditionally fabricated ones shown in Table 3.31. The Vickers hardness of the Al2O3 ceramic parts is summarized in order to check the practicability of SLA technique. Below 1600°C, the hardness increased with the sintering temperature. Samples sintered at 1600°C and 1650°C had a hardness of 17.7 and 17.9 GPa, respectively, which were the highest among all the sinterin g temperatures.

160

Additive and traditionally manufactured components

H

Layers

W L

Dimensional Shrinkage (%)

2.8

(A)

2.4 2.0 1.6 1.2 0.8 0.4 0.0 Height

Dimensional shrinkage (%)

24

Length

Width

Height Length Width

(B)

22 20 18 16 14 12 1400

1450

1500 1550 1600 Temperature (°C)

1650

Fig. 3.107 Shrinkage changes in Al2O3 sample after different processing steps (H—height, L—length, W—width): (A) debinded body and (B) sintered parts. (An, D, Li, H., Xie, Z., Zhu, T., Luo, X., Shen, Z., Ma, J., 1917. Int. J. Appl. Ceram. Tech. 14, 836. With kind permission of John Wiley and Sons.)

The variation in microhardness was strongly affected by the microstructure, such as porosity and grain size. As mentioned above, the samples were not dense until temperature raise to 1600°C.

3.3.6 Conventionally produced alumina It is puzzling that in the vast literature on alumina, almost no data on pure, undoped or unalloyed alumina exist. More so, since crystalline alumina occurs naturally as the mineral corundum and produced artificially because of its extensive use as a gemstone, abrasive—which makes it suitable for wear applications—electrical insulators and components of cutting tools to mention a few. Health and medical applications include the use as a material in hip replacements. Hardness is often a measured property in ceramics which is easier to perform than other deformation techniques. It generally measures the resistance of the material to permanent plastic deformation. Vickers hardness is usually

Testing: Comparison of AM data with traditionally fabricated

161

Fig. 3.108 The microstructure of Al2O3 ceramics sintered at various sintering temperatures: (A) 1200°C; (B) 1400°C; (C) 1500°C; (D) 1600°C; and (E and F) interlayered structure at 1600°C. (An, D, Li, H., Xie, Z., Zhu, T., Luo, X., Shen, Z., Ma, J., 1917. Int. J. Appl. Ceram. Tech. 14, 836. With kind permission of John Wiley and Sons.)

performed with a diamond pyramid indenter and requires the measurement of the diagonal for the calculation of the hardness value. The impression, namely the indent, is generally the applied load (P) dependent. The effect of the load is illustrated for slip cast (SC) alumina in Fig. 3.110. The variation of the indent with the load applied is summarized in Table 3.32 according to Eq. (3.46). In Fig. 3.110 each point represent the average of 30 values. The hardness values decrease with the increase in the indentation load until at high load the hardness becomes constant. A detailed description of hardness can be

162

Additive and traditionally manufactured components

Microhardness (GPa)

20

6

12

8

4 1400

1450

1500 1550 1600 Temperature (°C)

1650

Fig. 3.109 Hardness as a function of the sintering temperature. (An, D, Li, H., Xie, Z., Zhu, T., Luo, X., Shen, Z., Ma, J., 1917. Int. J. Appl. Ceram. Tech. 14, 836. With kind permission of John Wiley and Sons.) Table 3.31 A comparison of mechanical properties of Al2O3 parts via 3D printing with other traditional shaping methods. References

Shaping method

Hardness (GPa)

27 28 29 4

Tape casing Gel casting Slip casting Injection molding

15.91 10.66 15.1 >20

An, D., Li, H., Xie, Z., Zhu, T., Luo, X., Shen, Z., Ma, J., 1917. Int. J. Appl. Ceram. Tech. 14, 836. With kind permission of John Wiley and Sons.

found in mechanical properties of Ceramics (Pelleg, 2014). Thus reproducing Eqs. (1.26) and (1.27) below as H ¼α

P d2

(3.44)

and

2Psin ϕ =2 Load 1:854P ¼ DPH ¼ ¼ Hv ¼ 2 area d2 d

(3.45)

where ϕ ¼ 136 degrees is the indenter angle between the two opposite faces, DPH is the diamond pyramid hardness, i.e., the Vickers hardness,

163

Testing: Comparison of AM data with traditionally fabricated

3100 2800

HV

2500 2200 1900 1600 1300 0

5

10

15

20

25

30

35

40

45

50

F (N)

Fig. 3.110 Vickers hardness as a function of the applied load for SC-Al2O3 ceramics  c, L., Lalic, M., Šolic, S., 2009. Kovove Mater. (mean value of standard deviation). (Curkovi  c.) 47, 89. With kind permission of Dr. Curkovi Table 3.32 Average values of indents and standard deviation as a function of the applied load. Load (kg)

Average indent size (μm)

Standard deviation

0.05 0.1 0.2 0.5 1 3 5

5.7 8.6 13.3 21.9 34.0 61.2 78.6

0.23 0.33 0.43 0.78 0.91 1.47 1.7

  Curkovi c, L., Lalic, M., Sˇolic, S., 2009. Kovove Mater. 47, 89. With kind permission of Dr. Curkovi c.

Hv, d is the diagonal and P is the load. Clearly, the Vickers hardness is the load divided by surface area A (namely P/A) of the pyramid-shaped indentation which requires measuring of the diagonals. Meyer has proposed an empirical expression between the applied load and indentation size as P ¼ Ad n

(3.46)

Here A is a constant characteristic of the material and n is Meyer’s index, which is a measure of the indentation size effect. For ceramics n is between 1.5 and 2.0 and is obtained from experimental results such as illustrated for

164

Additive and traditionally manufactured components

1.7 1.4

y = 1.7461x + 3.5917 R2 = 0.9991

log F

1.1 0.8 0.5 0.2 -0.1 -0.4 -2.3

-2.1

-1.9

-1.5

-1.7 log d

-1.3

-1.1

Fig. 3.111 Vickers hardness data on SC-Al2O3 ceramics according to Meyer’s law.  c, L., Lalic, M., Šolic, S., 2009. Kovove Mater. 47, 89. With kind permission of (Curkovi  Dr. Curkovic.) Table 3.33 Regression analysis of experimental data according to Eq. (3.46). Sample

n

log A

A (N mm–n) Correlation factor R2

SC-Al2O3 1.7461  0.023 3.5917  0.039 3906

0.9991

 Curkovi c, L., Lalic, M., Sˇolic, S., 2009. Kovove Mater. 47, 89. Open access.

SC alumina in Fig. 3.111. Table 3.33 summarizes the results. An alternative analysis of the indentation size effect of Meyer is the proportional specimen resistance (PSR) of Li and Bradt (1993). An expression has been derived from experimental measurements as P ¼ α1 d + α2 d 2

(3.47)

Eq. (3.47) can be rewritten as P ¼ α1 + α2 d d

(3.48)

By the linear expression of P/d against d from Fig. 3.112 the parameters α1 and α2 can be evaluated. The parameter α1 is related to the proportional resistance of the specimen, while α2 represent the true hardness, which is load independent. Note that P is used in the equations, whereas in the curves

165

Testing: Comparison of AM data with traditionally fabricated

630 550

y = 7246.6x + 51.2 R2 = 0.9986

F/d (N/mm)

470 390 310 230 150 70 0.005

0.02

0.035

0.05 d (mm)

0.065

0.08

Fig. 3.112 Vickers hardness data on SC-Al2O3 ceramics according to the PSR model.  c, L., Lalic, M., Šolic, S., 2009. Kovove Mater. 47, 89. With kind permission of (Curkovi  c.) Dr. Curkovi

F has the same meaning as P. A modified PSR model—the MPSR—was suggested by Gong et al. (1999) which considers surface residual stress contribution to hardness induced by machining and expressed as P ¼ P0 + α1 d + α2 d2

(3.49)

The load vs indentation size according to the MPSR model is illustrated in Fig. 3.113. The results according to the PRS and MPRS models are listed in Tables 3.34 and 3.35, respectively. P0 is a constant related to the surface residual stresses associated with the specimen preparation (machining, polishing, etc.). 3.3.6.1 Temperature dependence The change in hardness with temperature for several nominally pure polycrystalline aluminas including single-crystal sapphire (α-Al2O3) was measured. As can be seen from Table 3.36 a small volume fraction of a second phase was present in the polycrystalline alumina. The second phase is an aluminosilicate as determined by TEM and SEM and the SEM micrograph is illustrated in Fig. 3.114. Our concern in this section relates to the nominally pure alumina (i.e., the 0.2 vol% material). As can be seen, the hardness of the single crystal and the polycrystalline aluminas have the same

166

Additive and traditionally manufactured components

48

y = 7454.2x2 + 34.408x + 0.1839 R2 = 0.9999

40

F (N)

32 24 16 8 0 0.005

0.02

0.035

0.05

0.065

0.08

d (mm)

Fig. 3.113 The applied load versus indentation size according to the modified PSR  c, L., Lalic, M., Šolic, S., 2009. Kovove Mater. 47, model for SC-Al2O3 ceramics. (Curkovi  c.) 89. With kind permission of Dr. Curkovi

Table 3.34 Regression analysis results of experimental data according to a PSR model. Sample

a1 (N mm21)

a2 (N mm22)

Correlation factor R2

SC-Al2O3

51.2  5.1

7246.6  123.2

0.9986

  Curkovi c, L., Lalic, M., Sˇolic, S., 2009. Kovove Mater. 47, 89. With kind permission of Dr. Curkovi c.

Table 3.35 Parameters Fo, α1, and α2 of the MPRS model according to Eq. (3.48) for SC-Al2O3 ceramics. Sample

F0 (N)

a2 (N mm21)

a2 (N mm22)

Correlation factor R2

SC-Al2O3

0.1839  0.229

34.408  16.078

7454.2  189.5

0.9999

  Curkovi c, L., Lalic, M., Sˇolic, S., 2009. Kovove Mater. 47, 89. With kind permission of Dr. Curkovi c.

167

Testing: Comparison of AM data with traditionally fabricated

Table 3.36 Properties of materials used. Material

Sapphire Vistalb AD999b AD96b AD90b

Second phase (vol%)

Grain size (μm)

μ (GPa)

H0 (GPa)

0.2 0.2 7 18

20 3 11 4

174 161 158 125 113

23.2 23.2 23.2 17.6 15.8

a

a

Adolph Meller Co. Providence, RI: surface orientation (2110). Coors Porcelain Co., Golden, CO. Alpert, C.P., Chan, H.M., Bennison, S.J., Lawn, B.R., 1988. J. Am. Ceram. Soc. 71, C 371. With kind permission of John Wiley and Sons.

b

Fig. 3.114 SEM of polished section of debased (AD90) alumina, showing pockets of aluminosilicate glass phase between the Al2O3 grains. Note that the latter grains make close contact with adjacent neighbors, i.e., the Al2O3 structure is connected. (Alpert, C.P., Chan, H.M., Bennison, S.J., Lawn, B.R., 1988. J. Am. Ceram. Soc. 71, C 371. With kind permission of John Wiley and Sons.)

hardness values. The hardness decreases with temperature increase as expected and is illustrated in Fig. 3.115. One of the empirical relations expressing the temperature dependence is given by H ¼ H0 ð1  T =T0 Þ

(3.50)

3.3.6.2 Hardness of coatings Decreased porosity in alumina coatings can be achieved by gas tunnel plasma spraying at atmospheric pressure. Clearly, alumina exhibiting porosity exhibits diminished strength properties which limits its applications. Higher

168

Additive and traditionally manufactured components

30 Sapphire Vistal AD999 AD96 AD90

25

Hardness (GPa)

20

15

10

5

0

0

400

800

1200

1600

Temperature (°C)

Fig. 3.115 Plot of hardness versus temperature for the alumina materials listed in Table 3.36. Solid lines are linear fits to Eq. (3.50). (Alpert, C.P., Chan, H.M., Bennison, S.J., Lawn, B.R., 1988. J. Am. Ceram. Soc. 71, C 371. With kind permission of John Wiley and Sons.)

quality alumina coatings can be obtained by the gas tunnel plasma spraying than with the conventional plasma spraying. In this method, the sprayed particles are in a fully molten state on contact with the substrate. The decreased porosity improves, in general, the mechanical properties and hardness among them. High Vickers hardness of alumina Hv ¼ 1200–1600 were obtained by the gas tunnel plasma spraying. A schematic view of the gas tunnel plasma spraying apparatus is shown in Fig. 3.116. The spraying distance, namely the distance between the torch and the substrate is critical as indicated by Fig. 3.117, where Lp designates the critical distance. As can be seen, it depends on the power applied. The coating is formed by scanning the substrate in front of the torch. The Vickers load was 300 g with a dwelling time of 15 s. In the figure average hardness of the top of the coating is indicated of at least 10 measurements. The Vickers hardness increases with decreasing spray distance as seen in Fig. 3.117. As seen in Fig. 3.117 the Vickers hardness at the surface of the coating is higher

169

Testing: Comparison of AM data with traditionally fabricated

Power supply

Vortex generator

L B

Gas divertor nozzle Powder feeder

Working gas (Ar)

Substrate

Inside nozzle electrode

Fig. 3.116 Schematic diagram of the gas tunnel plasma spraying apparatus, A is convention plasma torch. B is gas tunnel plasma torch. L is spraying distance. (Kobayashi, A., 1996. J. Thermal Spray Technology 5, 298. With kind permission of Springer Nature.)

Vickers hardness, Hv

1500

P = 30 kW 1000 Lp

P = 20 kW

500 0

20

30 40 50 60 70 Spraying distance, L (mm)

80

90

Fig. 3.117 Dependency of Vickers hardness of alumina coating on spraying distance at 20 and 30 kW P. Lp is critical spraying distance. (Kobayashi, A., 1996. J. Thermal Spray Technol. 5, 298. With kind permission of Springer Nature.)

170

Additive and traditionally manufactured components

P = 20 kW L = 30 mm

Vickers hardness, Hv

1500

1000

Surface

Substrate

500 0

100 200 300 400 Distance from coating surface, l (mm)

500

Fig. 3.118 Distribution of the Vickers hardness on the cross section of the alumina coating in the thickness direction. (Kobayashi, A., 1996. J. Thermal Spray Technol. 5, 298. With kind permission of Springer Nature.)

than that near the substrate, and the distribution of hardness shows a parabolic distribution. The peak Vickers hardness at a distance of 200 μm is >1300 Hv. The hardness also varies with the power input, reaching the highest value of 1500 at a spraying distance of 30 mm and a power of 30 kW. Very high hardness values can be obtained at a distance L < Lp. The distribution of the hardness shown in Fig. 3.118 where the hardness is plotted against the distance is a result of measurements performed on a coating of 450 μm. The distribution at a higher power input of 30 kW is illustrated in Fig. 3.119. The distribution is similar but the level of hardness values is higher. Clearly, protective coatings to be effective must have a microstructure without porosity and of high density, which can be achieved by gas tunnel plasma spraying. 3.3.6.3 Hardness of alumina films Of the many uses of Al2O3 films such as optical applications, insulating and protection layer, magnetic head recording its use in microelectronics, namely device fabrication, should be emphasized. Here we discuss the characteristics of alumina deposited by RF sputtering on Al2O3-TiC substrate. Atomic force microscope (AFM) was used for the characterization of the

171

Testing: Comparison of AM data with traditionally fabricated

P = 30 kW L = 30 mm

Vickers hardness, Hv

1500

1000

Surface

Substrate

500 0

100 200 300 400 Distance from coating surface, l (mm)

500

Fig. 3.119 Distribution of the Vickers hardness on the cross section of the alumina coating in the thickness direction in the case of large power input. (Kobayashi, A., 1996. J. Thermal Spray Technol. 5, 298. With kind permission of Springer Nature.)

film. The mechanical properties of the film were evaluated by hardness measurements. AFM was used to investigate the surface roughness of Al2O3 which is shown in Fig. 3.120. The scan size for all images was 5  5 μm2 and as seen in Fig. 3.120 the film roughness significantly increased with target sputtering power. Clearly, the film thickness also increases with the increase of power as shown in Fig. 3.121. The hardness (and elastic modulus) vs the target sputtering power is seen in Fig. 3.122. It increases with an increase in power. Table 3.37 summarizes the hardness values for the various sputtering power. Thus the thickness and hardness increase with the sputtering power, but this is on the expense of the surface roughness, which increases with the power. Thus we have discussed in addition to bulk alumina three additional aspects of importance regarding the applications of Al2O3. These are the temperature effect, which is important because alumina is intended for high-temperature use, the coating properties which is of interest in the application of coatings as an insulator and against corrosion and environmental effects and the film properties which is an important ingredient in microelectronic materials and also in its use as magnetic head recording.

172

Additive and traditionally manufactured components

nm

5.0

µm 4

0

5

3

-5.0

0

2 1

2.50 µm

5.00

(c) 5.0

nm

0

µm 5 4

-5.0

3

2

0

2.50 µm

1

5.00

(a) nm

5.0

0 µm 4 2

3

-5.0

0

1

2.50 µm

5.00

(b)

5.0

nm

0 5 4 2 1

3

µm -5.0

0

2.50 µm

5.00

Fig. 3.120 3D AFM image and cross section of Al2O3 films deposited at different target sputtering powers of (a) 8 kW, (b) 7 kW, (c) 6 kW, and (d) 5 kW. (Panitchakan, H., Limsuwan, P., 2012. Procedia Eng. 32, 902. With kind permission of Elsevier.)

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Testing: Comparison of AM data with traditionally fabricated

80.0

8.00

70.0

7.50

60.0

7.00

50.0

6.50

Er (GPa)

40.0

5.00

Hardness (GPa)

Elastic modulus ER (GPa)

Fig. 3.121 SEM micrograph of Al2O3 film deposited at target sputtering powers of (a) 6 kW and (b) 8 kW. (Panitchakan, H., Limsuwan, P., 2012. Procedia Eng. 32, 902. With kind permission of Elsevier.)

HD(GPa) 5.50

30.0

5.00

20.0 5

6

7

8

Target sputtering power (kW) Fig. 3.122 The plots of elastic modulus and hardness of Al2O3 film with target sputtering power. (Panitchakan, H., Limsuwan, P., 2012. Procedia Eng. 32, 902. With kind permission of Elsevier.)

174

Additive and traditionally manufactured components

Table 3.37 Hardness and elastic modulus of Al2O3 films deposited at various target sputtering powers. Sputtering power (kW)

Substrate bias voltage (V)

Elastic modulus (GPa)

Hardness (GPa)

5 6 7 8

150 150 150 150

27.6 50.3 61.9 62.8

6.6 7.2 7.5 7.6

Panitchakan, H., Limsuwan, P., 2012. Procedia Eng. 32, 902. With kind permission of Elsevier.

References Alcisto, J., Enriquez, A., Garcia, H., Hinkson, S., Steelman, T., Silverman, E., Valdovino, P., Gigerenzer, H., Foyos, J., Ogren, J., Dorey, J., Karg, K., McDonald, T., Es-Said, O.S., 2011. J. Mater. Eng. Perform. 20, 203. Amsterdam, E., Kool, G., 2009. 25th ICAF (International Committee on Aeronautical Fatigue) Symposium. Rotterdam, May 27–29, 1261–1274. Carroll, B.E., Dinda, G.P., Song, L., Mazumder, J., 2008. Metall. Mater. Trans. A 39, 2914. Chen, H.S., 1980. Rep. Prog. Phys. 43, 533. Cottam, R., Brandt, M., 2011. Phys. Procedia 12, 323. Dinda, G.P., Song, L., Mazumder, J., 2008. Metall. Mater. Trans. A 39, 2914. Feilden, E., Blanca, E.D.-T., Giuliani, F., Saiz, E., Vandeperre, L., 2016. J. Eur. Ceram. Soc. 36, 2525–2533. Gonzalez, J.A., Mireles, J., Lin, Y., Wicker, R.B., 2016. Ceram. Int. 42, 1–7. Gong, J., Wu, J., Guan, Z., 1999. J. Eur. Ceram. Soc. 19, 2025. Griffith, M.L., Ensz, M.T., Puskar, J.D., Robino, C.V., Brooks, J.A., Philliber, J.A., Smugeresky, J.E., Hofmeister, W.H., 2000. MRS Proceeding 625. . Huang, T., Mason, M.S., Zhao, X., Hilmas, G.E., Leu, M.C., 2009. Rapid Prototyp. J. 15, 88–95. Keicher, D.M., Miller, W.D., 1998. Met. Powder Rep. 53, 26. Kobryn, P.A., Semiatin, S.L., 2001. Proceedings of 12th Solid Freeform Fabrication Symposium. Austin 179. Larosa, M.A., Jardini, A.L., de Carvalho Zavaglia, C.A., Kharmandayan, P., Calderoni, D.R., Filho, R.M., 2015. Adv. Mech. Eng.. 6https://doi.org/10.1155/ 2014/945819. Lewis, G.K., Schlienger, E., 2000. Mater. Des. 21, 417. Li, H., Bradt, R.C., 1993. J. Mater. Sci. 28, 917. Li, J., Leu, M.C., Hilmas, G.E., 2015a. Bourell, D., Beaman, J., Crawford, R., Fish, S., Marcus, H., Seepersad, C. (Eds.), Solid Freedom Fabrication Symposium. Austin, TX, USA, pp. 319–331. Li, X., Reynolds, A.P., Baoqiang, C., Jialuo, D., Williams, S., 2015b. TMS Annual Spring Meeting Supplemental Proceeding, TMS 2015. p. 445. With kind permission of Springer Nature. Liu, Z.H., Nolte, J.J., Packard, J.I., Hilmas, G., Dogan, F., Leu, M.C., 2007. Hinduja, S., Fan, K.C. (Eds.), Proceedings of the 35th International MATADOR Conference. In: 3, Springer, Taipei, Taiwan, pp. 351–354. Maleksaeedi, S., Eng, H., Wiria, F.E., Ha, T.M.H., He, Z., 2014. J. Mater. Process. Technol. 214, 1301–1306. Palmer, T.A., Beese, A.M., 2015. Acta Mater. 87, 309.

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Pelleg, J., 2012. Mechanical Properties of Materials. Springer. Pelleg, J., 2014. Mechanical Properties of Ceramics. Springer. Qiu, C., Ravi, G.A., Dance, C., Ranson, A., Dilworth, S., Attallah, M.M., 2015. J. Alloys Compd. 629, 351. Rønneberg, T., 2016. Characterization of Aluminium Components Produced by Additive Manufacturing. Norwegian University of Science and Technology, Trondheim. Schwentenwein, M., Homa, J., 2015. Int. J. Appl. Ceram. Technol. 12, 1–7. Tong, J., Bowen, C.R., Persson, J., Plummer, A., 2017. Mater. Sci. Technol. 33, 138. Wang, Z., Palmer, T.A., Beese, A.M., 2016. Acta Mater. 110, 226. Yadroitsev, I., Krakhmalev, P., Yadroitsava, I., 2014. J. Alloys Compd. 583, 404. Yao, B., Ma, X.-L., Lin, F., Ge, W.-J., 2015. Rare Met. 34, 445. Yu, J., Rombouts, M., Maes, G., Motmans, F., 2012. Phys. Procedia 39, 416. Zhai, Y., Galarraga, H., Lados, D.A., 2015. Procedia Eng. 114, 658. Zhang, X.D., Zhang, H., Grylls, R.J., Lienert, T.J., Brice, C., Fraser, H.L., Keicher, D.M., Schlienger, M.E., 2001. J. Adv. Mater. 33, 17. Zhang, S., Lin, X., Chen, J., Huang, W., 2009. Rare Met. 28, 537.

Further reading Abd-Elghany, K., Bourell, D.L., 2012. Rapid Prototyp. J. 18, 420. Accuratus Ceramic Corporation, 2013. Phillipsburg, New Jersey 08865.G115. An, D., Li, H., Xie, Z., Zhu, T., Luo, X., Shen, Z., Ma, J., 1917. Int. J. Appl. Ceram. Technol. 14, 836. Arulselvan, M., Ganesan, G., 2013. Int. J. Recent Technol. Eng. 2, 47. Atlas Steels, 2013. Atlas Steels. Technical Department. Revised: August. Baldenebro-Lopez, F.J., Gomez-Esparza, C.D., Corral-Higuera, R., Arredondo-Rea, S.P., Pellegrini-Cervantes, M.J., Ledezma-Sillas, J.E., Martinez-Sanchez, R., HerreraRamirez, J.M., 2015. Materials 8, 451. Beese, A.M., Carroll, B.E., 2016. JOM 68, 724. Clinton Aluminum of the Aluminum Association, n.d. MatWeb, Your Source for Materials Information, List of Data  Curkovi c, L., Lalic, M., Sˇolic, S., 2009. Kovove Mater. 47, 89. Dutta, B., Froes, F.H., 2015. The additive manufacturing (AM) of titanium alloys. Chapter 24, In: Sci. Tech. Appl, p. 447. Ga´lvez, F., Rodrı´guez, J., Sa´nchez Ga´lvez, V., 2000. J. Phys. IV France. 10. Pr9-323. Ghazanfari, A., Li, W., Leu, M., Watts, J., Hilmas, G., 2017. Int. J. Appl. Ceram. Technol. 14, 486. Gonzalez, J.A., Mireles, J., Lin, Y., Wicker, R.B., 2016. Ceram. Int. 42, 10559. Goodfellow, 2008–2020. Glass and Ceramics for Science and Industry, Properties of Alumina. Gray III, G.T., Livescu, V., Rigg, P.A., Trujillo, C.P., Cady, C.M., Chen, S.R., Carpenter, J.S., Lienert, T.J., Fensin, S., 2015. EPJ Web of Conferences 94. 02006. Kobayashi, A., 1996. J. Therm. Spray Technol. 5, 298. Longhitano, G.A., Larosa, M.A., Jardini, A.L., de Carvalho Zavaglia, C.A., Ierardi, M.C.F., 2018. J. Mater. Process. Technol. 252, 202. Ma, Y., Cuiuri, D., Hoye, N., Li, H., Pan, Z., 2014. J. Mater. Res. 29, 2066. Moreira, P.M.G.P., Santos, T., Tavares, S.M.O., Richter-Trummer, V., Vilac¸a, P., de Castro, P.M.S.T., 2009. Mater. Des. 30, 180. Panitchakan, H., Limsuwan, P., 2012. Procedia Eng. 32, 902. Poondla, N., Srivatsan, T.S., Patnaik, A., Petraroli, M., 2009. J. Alloys Compd. 486, 162. Qu, S., Huang, C.X., Gao, Y.L., Yang, G., Wu, S.D., Zang, Q.S., Zhang, Z.F., 2008. Mater. Sci. Eng. A 475, 207.

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Rao, P.N., Singh, D., Jayaganthan, R., 2013. Mater. Sci. Technol. 29, 76. Reichardt, A., Dillon, R.P., Borgonia, J.P., Shapiro, A.A., McEnerney, B.W., Momose, T., Hosemann, P., 2016. Mater. Des. 104, 404. Savchenko, N., Sevostyanova, I., Sablina, T., G€ omze, L., Kulkov, S., 2014. AIP Conf. Proc. 1623, 547. Schmid, F., Harris, D.C., 1998. J. Am. Ceram. Soc. 81, 885. Semiatin, S.L., Seetharaman, V., Weiss, I., 1999. Mater. Sci. Eng. A 263, 257. Shang, S.-Z., Lu, G.-M., Tang, X.-L., Zhao, Z.-X., Wu, C.-m., 2010. Trans. Nonferrous Met. Soc. China 20, 1725. Song, B., Nishida1, E., Sanborn, B., Maguire, M., Adams, D., Carroll, J., Wise, J., Reedlunn, B., Bishop, J., Palmer, T., 2017. J. Dyn. Behav. Mater. 3, 412. Terada, D., Kaneda, Y., Horita, Z., Matsuda, K., Hirosawa, S., Tsuji, N., 2014. IOP Conf. Series: Mater. Sci. Eng. 63, 012088. Uddin, S.Z., Murr, L.E., Terrazas, C.A., Morton, P., Roberson, D.A., Wicker, R.B., 2018. Addit. Manuf. 22, 405.

CHAPTER FOUR

Dislocations in AM and traditional manufacturing: A comparison

4.1 Introduction A detailed discussion on dislocations can be found in Mechanical Properties of Materials (Pelleg, 2012) and Mechanical Properties of Ceramics (Pelleg, 2014). Briefly, mechanical properties of materials are determined by defects, particularly those known as dislocations. The postulate for the existence of such defects by the fathers of the dislocation theory, namely, Taylor, Orowan, and Polanyi, has been confirmed by numerous research studies. The actual existence of dislocations has been observed by TEM, Field Ion Microscopy (FIM), and atom probe techniques (permits direct observation of dislocations on an atomic scale). An etch pit is an indirect technique to detect the presence of dislocations in solids. Therefore, in this chapter— essential to understand plastic deformation to be discussed in the next chapter—the emphasis is on the comparison of the dislocation structures in selected alloys and ceramics. As mentioned in the preface, the alloys to be considered are Ti-6Al-4V, Al 6061, steel 304L, and alumina representing ceramics.

4.1.1 In AM Ti-6Al-4V In a recent review on AM using electron beam melting (EBM) to fabricate materials, among them Ti-based alloys is discussed. Due to their light weight and high specific strength, Ti-based alloys such as Ti-6Al-4V have a wide range of applications in aeronautic systems. Further, the Ti-6Al-4V alloy has been a popular orthopedic joint replacement. The presence of Al and V increases the α (HCP) to β (BCC) phase transition temperature from 885°C to 995°C and the β-phase form under strain martensite (HCP α0 ). This martensitic phase—known also as an acicular phase—is shown in Fig. 4.1 as an optical micrograph where a TEM image showing high-density dislocations in the α-phase is inserted. The vertical section shown is parallel Additive and Traditionally Manufactured Components https://doi.org/10.1016/B978-0-12-821918-8.00004-8

© 2020 Elsevier Inc. All rights reserved.

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Fig. 4.1 Vertical section view (optical micrograph) for an EBM-fabricated Ti-6Al-4V cylindrical component showing acicular α-phase grains surrounded by interfacial β-phase (black). The TM image insert shows high dislocation density in α-phase grains. (Murr, L.E., Gaytan, S.M., Ramirez, D.A., Martinez, E., Hernandez, J., Amato, K.N. et al., 2012. J. Mater. Sci. Technol. 28, 1. With kind permission of Elsevier.)

to the build direction. The probable reason for the dislocations seen in the α-phase is the consequence of the thermal stresses induced by solidification. The solidification rate depends on the volume. An additional dislocation structure is seen in Fig. 4.2. These figures are a part of figures of the review of fabrication by additive manufacturing of various alloys. As mentioned, an important application of Ti-6Al-4V is in orthopedic and dental bone implants due to their long-term load-bearing potential. Their use is a consequence of good biocompatibility, high strength-to-weight ratio, fatigue resistance, excellent corrosion resistance, and lower modulus compared with other biomaterials. In further considering dislocations in α0 martensitic grains, Fig. 4.3 provides additional information on this subject. Microstructures of struts in as-built Ti-6Al-4V are illustrated in the figure,

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Fig. 4.2 TEM bright-field image showing a fine α-phase microstructure for an EBM- fabricated Ti-6Al-4V. Note dislocation structure in α-grains. (Murr, L.E., Gaytan, S.M., Ramirez, D.A., Martinez, E., Hernandez, J., Amato, K.N. et al., 2012. J. Mater. Sci. Technol. 28, 1. With kind permission of Elsevier.)

including dislocations in the martensitic α0 - phase. The heat treatment of the as-built samples at 650°C for 4 h and then cooling down in the furnace under Ar atmosphere to room temperature for stress relieving and homogenization of the structure is seen in this figure. The heat treatment resulted in the micrographs shown in Fig. 4.4. The dislocation density in the α-phase of the heattreated Ti-6Al-4V is lower than in the α0 as-built structure. This is seen in Fig. 4.4d. It is worth noting that TPMS in differential geometry is a minimal surface in R3 that is invariant under a rank-3 lattice of translations. These faces have the symmetries of a crystallographic group. The dislocations substructure for a transverse section cut from a cylindrical build in the α-phase structure observed by TEM bright-field is seen in Fig. 4.5. The dislocation density in Fig. 4.5b is 5  109/cm2. The differences in the α-phase microstructures generally are associated with the residual hardness variations and the hardness is related to the yield stress according to an empirical relation of.

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Fig. 4.3 Microstructure of the struts in as-built Ti-6Al-4V TPMS structures. (a) Optical, (b) SEM micrographs, (c) TEM bright-field image with the inset of corresponding selected-area diffraction pattern, and (d) high magnification TEM micrograph showing dislocations and crystallographic defects (either stacking or twin faults) in the α0 martensite grains. TPMS stands for triply periodic minimal surface. (Yan, C., Hao, L., Hussein, A., Young, P., 2015. J. Mech. Behav. Med. Mater. 51, 61.)

Yield stress ffi Hv=3

(4.1)

In Ti-6Al-4V, for example, obtained by single-melt-pass EBM, the yield stress varied from 1 to 1.2 GPa with the corresponding microhardness Hv variation from 3.5 to more than 4.1 GPa. In addition to the acicular α-phase microstructure changes, variation in the cooling rate or solidification rate can also cause differences in the dislocation density within the α-phase. These features are illustrated in Figs. 4.6 and 4.7. An illustrative example of Ti-6Al-4V as an implant is presented in Fig. 4.8 where a knee plant is shown. Note: The knee is the largest weight-bearing joint of the body. Three bones meet to form the knee joint: the femur (thighbone), tibia (shinbone), and patella (kneecap). Ligaments and tendons act like strong ropes to hold the bones together. They also work as restraints allowing some types of knee movements, and not

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Fig. 4.4 Microstructure of the struts in heat-treated Ti-6Al-4V TPMS structures. (a) Optical and (b) SEM micrographs, (c) TEM bright-field image with the inset of corresponding selected-area diffraction pattern, and (d) high magnification TEM micrograph showing dislocations. TPMS stands for triply periodic minimal surface. (Yan, C., Hao, L., Hussein, A., Young, P., 2015. J. Mech. Behav. Med. Mater. 51, 61.)

Fig. 4.5 Examples of TEM bright-field images of dislocation substructure in the α-phase of the cylindrical Ti-6Al-4V build: Transverse section views. (a) α-phase grains and β (black) boundary transition zones. (b) Magnified view showing dislocation substructure in α. (Murr, L.E., Quinones, S.A., Gaytan, S.M., Lopez, M.I., Rodela, A., Martinez, E.Y. et al., 2009. J. Mech. Behav. Biomed. Mater. 2, 20. With kind permission of Elsevier.)

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Fig. 4.6 TEM bright-field image comparisons for single-melt-pass EBM fabrication of fully dense Ti-6Al-4V monoliths with variations in build thermal history to create dislocation density variations in the α-phase. (a) High hardness (HV ¼ 3.9 GPa), dislocation density approximately 1010 cm2. (b) Magnified view of dislocations in α-phase in (a). (c) Low hardness (HV ¼ 3.5 GPa), dislocation density approximately 107 cm2. (d) Magnified view of region in (c). (a) and (b) Scale bars, 1 mm. (Murr, L.E., Gaytan, S.M., Medina, F., Lopez, H., Martinez, E., Machado, B.I. et al., 2010. Phil. Trans. R. Soc. A 368, 1999; Open access. With kind permission of Elsevier.)

others. In addition, the way the ends of the bones are shaped help to keep the knee properly aligned. See the illustration below in Fig. 4.9 illustrating a proximal tibia fracture. The proximal tibia is the upper portion of the bone where it widens to help form the knee joint.

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Fig. 4.7 TEM bright-field image showing microstructural details for optimized EBM Ti-6Al-4V products. Dislocation-ledge structure along phase boundary. The selectedarea electron diffraction (SAED) pattern inset shows the hexagonal close-packed crystal orientation. Scale bar, 0.5 mm. (Murr, L.E., Gaytan, S.M., Medina, F., Lopez, H., Martinez, E., Machado, B.I. et al., 2010. Phil. Trans. R. Soc. A 368, 1999; Open access. With kind permission of Elsevier.)

4.1.2 In traditionally fabricated Ti-6Al-4V TEM bright-field images of the microstructure in the wrought Ti-6Al-4V sample is illustrated in Fig. 4.10 showing a dislocation network lying in the basal (0001) HCP α-phase. The dislocation density in the image of Fig. 4.10a representing the α-phase is 4  109/cm2 while in Fig. 4.10b a density of 1010/cm2 is indicated, corresponding to the α-phase region. Twins were also observed in the microstructure by TEM which contribute to the mechanical behavior of the wrought Ti-6Al-4V. However, the most effective strength and hardness contributions arise from the density of the dislocations in the microstructure. Dislocation arrays in an α-plate are seen in Fig. 4.11. Microstructural details in the wrought samples (Fig. 4.11a) are seen, including α-plates and β in Fig. 4.11b and dislocation substructures in Fig. 4.11c, where the corresponding dislocation density

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Fig. 4.8 Knee implant (tibial stem) prototype development and EBM processing. X-ray image for female right knee (total knee) replacement (femur, F; tibia, T). (Murr, L.E., Gaytan, S.M., Medina, F., Lopez, H., Martinez, E., Machado, B.I. et al., 2010. Phil. Trans. R. Soc. A 368, 1999; Open access. With kind permission of Elsevier.)

Femur (thighbone)

Patella Proximal tibia

Lateral collateral ligament

Cruciate ligaments

Tibial plateau

Fibula Tibia (shinbone)

Medial collateral ligament

Fig. 4.9 A proximal tibia fracture. (From The Body Almanac, American Academy of Orthopedic Surgeons, 2003.)

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Fig. 4.10 TEM (bright-field) images showing variations in dislocation substructures in α-phase regions of wrought Ti-6Al-4V. (a) Hexagonal dislocation network structure in tilted (0001) phase orientation. (b) Heavy dislocation substructure within α-phase regions. (Murr, L.E., Quinones, S.A., Gaytan, S.M., Lopez, M.I., Rodela, A., Martinez, E.Y. et al., 2009. J. Mech. Behav. Biomed. Mater. 2, 20. With kind permission of Elsevier.)

was 2  109 cm2. This is in contrast to 5  109–1010 cm2 measured for the EBM-1 sample illustrated in Fig. 4.10b. The experimental observations have indicated for the case of Ti-6Al-4V that specimens manufactured by EBM (an AM process) have strength and elongation comparable to sound wrought products. However, for direct comparison of the mechanical properties, the products in both the AM and the traditionally fabricated have to be by the best achievable production techniques under equivalent conditions of the starting material and its characteristics. Moreover, it has been also reported that in the EBM technique greater elongation can be obtained in the range of roughly 23%–92%, greater than the average elongation for high-strength Ti-6Al-4V forgings.

4.1.3 Motion of dislocations By in situ TEM deformation of Ti-6Al-4V at room temperature, dislocation motion was inspected. All dislocations have an a-type Burgers vector and glide essentially in prismatic or basal planes. At first, the dislocations are emitted from αS/β interfaces and then they move in preferential orientations along their screw direction. The motion of screw dislocations controls the strain rate. The Ti-6Al-4V is a two-phase alloy where the α-phase is HCP and the β-phase is BCC. The thermomechanical treatment determines the microstructure, which in turn determines the mechanical properties. The microstructure could be fully nodular, fully lamellar, or duplex with nodules and lamellar colonies. The properties of the Ti-6Al-4V are more complex because the presence of two phases and the α-phase can have

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Fig. 4.11 Wrought (W-1) Ti-6Al-4V microstructures. (a) Optical metallographic overview. (b) TEM image of acicular α with β boundaries corresponding to an area featured at the arrow in (a). (c) Magnified view of planar dislocation arrays in α-plate. (Murr, L.E., Esquivel, E.V., Quinones, S.A., Gaytan, S.M., Lopez, M.I., Martinez, E.Y. et al., 2009. Mater. Charact. 60, 96. With kind permission of Elsevier.)

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Fig. 4.12 Emission of a dislocation loop from an αS/β interface during an in situ experiment. The edge segment velocity is much larger than the screw segment one. (Castany, P., Pettinari-Sturmel, F., Crestou, J., Douin, J., Coujou, A., 2007. Acta Mater. 55, 6284. With kind permission of Elsevier.)

different chemical composition and different dislocation behavior. In Fig. 4.12, the sequence and emission of a dislocation loop is illustrated in an αS plate. The αS designates the α colonies as seen in Fig. 4.13. As a matter of fact, the lamellar colonies of αS/β are well illustrated in Fig. 4.14. Returning to Fig. 4.12, the dislocations are emitted from the interface of the αS- and β-phases at point A of Fig. 4.12a. The emitted dislocation expands as seen in images Fig. 4.12b and c. A preferential orientation along

Fig. 4.13 Detailed view of a lamellar colony with αS and β plates. In A and B, the β plates are discontinuous. (Castany, P., Pettinari-Sturmel, F., Crestou, J., Douin, J., Coujou, A., 2007. Acta Mater. 55, 6284. With kind permission of Elsevier.)

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Fig. 4.14 Microstructure of the investigated alloy with primary alpha nodules α and lamellar colonies αS/β. (Castany, P., Pettinari-Sturmel, F., Crestou, J., Douin, J., Coujou, A., 2007. Acta Mater. 55, 6284. With kind permission of Elsevier.)

the screw direction results in a large density of rectilinear screw segments, but departure from the screw orientation can occur close to the edge segment because of pinning points where bowing occurs as indicated by point B in the image of Fig. 4.12d. In this micrograph, the dislocation loop emerges at the sample surface at point C and the slip plane trace is visible. This dislocation glides in the basal plane and has an a-type Burgers vector (as the known invisibility criterion is determined by g∙ b ¼ 0). All gliding dislocations in this lamellar colony have the same Burgers vector, aligned in the screw direction and glide in the basal plane. Screw segments have a jerky motion, with very fast jumps. The dislocation motion can be inhibited by the formation of intrinsic obstacles. These obstacles are macrokinks formed by screw segment jumps. The sequence is illustrated in Fig. 4.15 where the motion of the screw segment is from the left to the right in the basal plane. In Fig. 4.15a a macrokink is created by the jump of a part of the dislocation. This is followed after about 13 s (during this time the dislocation is immobile) by the jump of the macrokink into another position as seen in Fig. 4.15b. The same event occurs between Fig. 4.15b and c.

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Fig. 4.15 Typical motion of a screw dislocation with the propagation of a macrokink (arrowed) along the same dislocation. The dotted line is at exactly the same position in every frame and is parallel to the screw direction. (Castany, P., Pettinari-Sturmel, F., Crestou, J., Douin, J., Coujou, A., 2007. Acta Mater. 55, 6284. With kind permission of Elsevier.)

Between Fig. 4.15c and d, the macrokink does not jump but glides slowly along the dislocation line. The slow continuous motion is along the screw direction. The screw direction is outlined by the dotted line. In Fig. 4.16, the creation of a dislocation loop is shown. Pinning of dislocation causes the

Fig. 4.16 Creation and expansion of a dislocation loop by the open-loop mechanism. At the end (e), the initial dislocation has moved away and has left two new mobile screw segments (arrowed). (Castany, P., Pettinari-Sturmel, F., Crestou, J., Douin, J., Coujou, A., 2007. Acta Mater. 55, 6284. With kind permission of Elsevier.)

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dislocation to bow out as seen in Fig. 4.16a, after which the dislocation slipped, returned to the screw character, and left behind a dislocation loop (see Fig. 4.16b). The loop expands under stress (Fig. 4.16c and d) and the initial dislocation moves away leaving behind two screw segments as illustrated in Fig. 4.16e. The loop is formed by a cross-slip (by the open-loop mechanism). This mechanism is the main dislocation multiplication mechanism. Cross-slip occurs frequently in this alloy and an example of the trace left by the propagating dislocations is seen in Fig. 4.17. The dislocation glides mainly in the basal plane and cross-slip occurs in a first-order pyramidal plane. The intrinsic obstacles of the screw dislocations are due to high lattice friction resulting from their three-dimensional core structure. Their core structure acts as intrinsic obstacles. Extrinsic obstacles also pin dislocation. In summary, dislocations observed in in situ TEM deformation experiments indicate that the gliding dislocations have an a-type Burgers vector and they are aligned preferentially with screw direction resulting in a large density of long rectilinear screw dislocations. The core structure causes jerky motion of the screw dislocations, cross- slip, and the formation of intrinsic

Fig. 4.17 Observation of a slip plane trace of a dislocation gliding mainly in the basal plane. The occurrence of cross-slip is visible (arrow) and takes place probably in a firstorder pyramidal plane. A first-order pyramidal plane trace (tr111) and the basal plane trace (trB) are reported. (Castany, P., Pettinari-Sturmel, F., Crestou, J., Douin, J., Coujou, A., 2007. Acta Mater. 55, 6284. With kind permission of Elsevier.)

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Fig. 4.18 A postmortem image of the same alloy deformed macroscopically to 0.3% plastic. The dislocation configuration is the same as in our in situ experiments: a dislocation loop, denoted by 1, emerges from an αS/β interface and the dislocations are rectilinear and aligned with their screw direction, like the dislocations denoted by 2 and 3. (Castany, P., Pettinari-Sturmel, F., Crestou, J., Douin, J., Coujou, A., 2007. Acta Mater. 55, 6284. With kind permission of Elsevier.)

obstacles including the creation of macrokinks and dislocation multiplication by the open-loop mechanism. The dislocations are omitted from the αS/β interface, glide in prismatic and basal planes, but basal plane glide is favored in lamellar colonies (Fig. 4.18).

4.2 Introduction AA6061 AA6061 is one of the most common aluminum alloys for general purposes containing magnesium and silicon. Aluminum alloys have been increasingly applied as a structural material in composite materials using a metal matrix due to their excellent mechanical properties and low weight. Al alloys are used in various industries, among them marine, defense, automotive transportation, and aerospace applications. The alloy is a precipitation hardening material, but it undergoes work hardening also. These two aspects are important strengthening methods, but when both are combined to improve the alloy properties, the strengthening mechanism is very impressive. The alloy is available as several grades such as the pretempered annealed grade known as 6061-O, tempered alloy designated as 6061-T6, which is solutionized and artificially aged, and the 6061-T651 alloy, which is solutionized, stress-relieved (stretched), and artificially aged. The mechanical properties of 6061 depend greatly on the temper or heat treatment.

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Alloy 6061 has an intermediate strength but it possesses excellent corrosion resistance and is readily weldable (not all Al alloys are easy to weld). Due to its resistance to corrosion, these alloys are used typically in marine structures, tank cars, pipelines, etc. It is important to realize that the alloy is available commercially in a variety of forms, shapes, sizes, etc.

4.2.1 AM of AA6061 Al alloy Only limited, if at all, illustrative examples on additive manufactured 6061 Al alloy is recorded in the literature despite its wide use in industrial applications. Therefore, instead of 6061, an Al alloy similar to that of 5083H116 is presented as an example for the dislocation structure. The similar behavior of these two alloys, AA 6061 and 5083-H116, can be seen from Figs. 4.19 and 4.20 illustrating their elastic moduli and the ultimate tensile stress. They are almost the same, in particular at the higher temperatures of 250–300°C. Recall that the 6061-T651 alloy is a solutionized, stress-relieved (stretched), and artificially aged Al alloy as indicated in the introduction to this section. The as-received dislocation cell (subgrain) structure is shown in Fig. 4.21.

70

Young’s Modulus (GPa)

60 50 40 30 5083-H116 6061-T651 Eurocode g [6] 5083-H321 [38] 6061-T651 [38]

20 10 0 0

100

200 300 Temperature (°C)

400

500

Fig. 4.19 5083-H116 and 6061-T651 elevated temperature Young’s modulus. (Summers, P.T., Chen, Y., Rippe, C.M., Allen, B., Mouritz, A.P., Case, S.W. et al., 2015. Fire Sci. Rev. 4, 3; Open access.)

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400 5083-H116 6061-T651 5083-H321 [38] 6061-T651 [38]

Ultimate Tensile Strength (MPa)

350 300 250 200 150 100 50 0 0

100

200 300 Temperature (°C)

400

500

Fig. 4.20 5083-H116 and 6061-T651 elevated temperature ultimate strengths. (Summers, P.T., Chen, Y., Rippe, C.M., Allen, B., Mouritz, A.P., Case, S.W. et al., 2015. Fire Sci. Rev. 4, 3; Open access.)

4.2.2 Dislocations in conventionally produced Al AA6061 The dislocations density and the subgrain (low angle) size in the base metal (BM) and a friction stir weld (FSW) 6061-T6 aluminum alloy were measured. The former was measured by XRD and the latter by neutron diffraction resulting in the dislocation densities of 4.5  1014 m2 and

Fig. 4.21 5083-H116 dislocation cell (subgrain) structure in the as-received state. (Summers, P.T., Chen, Y., Rippe, C.M., Allen, B., Mouritz, A.P., Case, S.W. et al., 2015. Fire Sci. Rev. 4, 3; Open access.)

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Fig. 4.22 Microscopy of the 6061-T6 Al alloy: TEM bright-field images of (a) base material and (b) FSW. The black line segments and dots indicate dislocations and precipitates in the TEM grain structure, respectively. (Woo, W., Ungár, T., Feng, Z., Kenik, E., Clausen, B., 2010. Metall. Mater. Trans. A 4, 1210.)

3.2  1015 m2, respectively. The subgrain sizes evaluated are 200 nm for BM and 160 nm for the FSW alloy. The significant increase in the dislocation density is likely to be due to severe plastic deformation during FSW. However, the comparison shows that the as-received Al rolled plate contains a larger number of dislocations (black line segments) and FSW (stir zone) has less dislocation as illustrated in Figs. 4.22a and b, respectively. The lower dislocation density in FSW is due to dynamic recrystallization (thermomechanically deformed). TEM bright-field images show the dislocation structure embedded grains of the base material (Fig. 4.22a) and the thermomechanically deformed stir zone of FSW (Fig. 4.22b). Ultrafine grained Al 6061 can be fabricated with reduced precipitation rate where the coarsening of recrystallized grains is inhibited and more dislocations are retained in this parent alloy. The microstructure after solid solution treatment and rolling indicates an average grain size of 700 nm, ill-defined grain boundaries, the absence of precipitates, and high dislocation density as illustrated in Fig. 4.23. In friction stir welding in the stir zone (SZ) region, equiaxed grains smaller than 2 μm with well-defined grain boundaries and lower dislocation density were observed (Fig. 4.24a) compared with the conditions represented in Fig. 4.23. This can be seen in TEM images in Fig. 4.24. The precipitate in SZ is 200 nm. In HAZ, the grains are somewhat smaller (Fig. 4.24d) and exhibit higher dislocation density (Fig. 4.24c). TEM images of specimens in water-cooled conditions are shown in Fig. 4.25. Again, in the HAZ condition, finer grains and larger dislocation density are obtained as illustrated in Fig. 4.25c. In addition, needle-like precipitates were observed

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195

500 nm

Fig. 4.23 TEM image of UFG 6061 Al alloy. (Liu, C.Y., Qu, B., Xue, P., Ma, Z.Y., Luo, K., Ma, M. Z. et al., 2018. J. Mater. Sci. Technol. 34, 112.)

Fig. 4.24 TEM images of (a) and (b) SZ and (c) and (d) HAZ of 800-Air. HAZ stands for heat-affected zone. (Liu, C.Y., Qu, B., Xue, P., Ma, Z.Y., Luo, K., Ma, M.Z. et al., 2018. J. Mater. Sci. Technol. 34, 112.)

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Fig. 4.25 TEM images of (a and b) SZ and (c and d) HAZ of 800-Water. HAZ stands for heat-affected zone. (Liu, C.Y., Qu, B., Xue, P., Ma, Z.Y., Luo, K., Ma, M.Z. et al., 2018. J. Mater. Sci. Technol. 34, 112.)

as illustrated in Fig. 4.25d. In the SZ specimens, equiaxed grains with sizes 1 μm and equiaxed rod-like precipitates of nanometers were obtained (Fig. 4.25a and b). At a temperature of 400°C (water temperature) the low heat input cannot lead to coarsening of recrystallized grains as indicated in the TEM image of Fig. 4.26a and b for SZ and Fig. 4.26c and d for HAZ. Fabrication at 400 water of the SZ samples providing a low heat input plays an important role in the hardness of FSW joints. AA 6061 is a precipitation hardened weldable Al alloy. Above, ultrafine grained fabrication of 6061 was discussed which can be produced by severe plastic deformation (SPD) and FSW. The fabrication methods considered in this section are: (1) rolling, characterized by nonequilibrium grain boundaries, and high dislocation density, no precipitates with good mechanical properties; (2) FSW in water, which inhibits grain coarsening in recrystallization, retains high density of dislocations and with reduced rate of precipitation. The alloys can be fabricated by cold rolling, and (3) specimens

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Fig. 4.26 TEM images of (a and b) SZ and (c and d) HAZ of 400-Water. Arrow in (d) denotes the precipitate. HAZ stands for heat-affected zone. (Liu, C.Y., Qu, B., Xue, P., Ma, Z.Y., Luo, K., Ma, M.Z. et al., 2018. J. Mater. Sci. Technol. 34, 112.)

welded in water that exhibited higher hardness and strength than those welded in air. The strengthening of the alloy fabricated in 400-water is a consequence of dislocations.

4.2.2.1 Pinning of dislocations in 6061 The mechanical properties of an alloy—in the present case of 6061—can be improved by dislocation pinning of particles. Nanoprecipitates obtained by warm laser shock peening (WLSP) technique, which is a new high strain rate strengthening process, can effectively pin dislocations. It is a thermomechanical process with great potential to improve also the fatigue life and surface hardness (Fig. 4.27). Surface hardness is an important factor per se but it also increases fatigue resistance by preventing initial crack formation. One of the important characteristics of WLSP is that high-density uniformly distributed dislocation

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350 300 400-Water

Stress (MPa)

250 200

800-Water

150 100

800-Air

50 0 0

2

4

6 8 Strain (%)

10

12

14

Fig. 4.27 Stress-strain curves of various FSW samples. (Liu, C.Y., Qu, B., Xue, P., Ma, Z.Y., Luo, K., Ma, M.Z. et al., 2018. J. Mater. Sci. Technol. 34, 112.)

structures and nanoscale precipitates are generated. The precipitates are formed by what is termed dynamic precipitation (DP) and serve as pinning agents for dislocation motion. Before laser shock peening, AA6061 samples are solutionized at 550°C for 3 h and subsequently quenched in water. A Q-switched ND-YAG laser, operating at a wavelength of 1064 nm with a pulse width [full-width at half-maximum (FWHM)] of 5 ns is used to deliver the pulsed laser. The laser power intensity I0 is adjusted by the Q-switched delay time and is calculated as 0:1E I0 ¼
 π ðd=2Þ2 t

(4.2)

where E is the pulse laser energy, d is the laser beam diameter, and t is the pulse width. The laser power intensity is one of the most important laser parameters, because it determines the peak pressure of the laser-induced plasma. Transmission electron microscopes are used to characterize the microstructures after WLSP. To prepare the TEM sample, a thin layer of the sample surface after shock peening is sectioned by a diamond saw. The sample is then polished to a thickness of 30 μm followed by thinning for the TEM observation. TEM images of AA6061 are shown in Fig. 4.28. The pinning effect depends on the distribution of the nanoprecipitates, its volume density, and size. The volume density of the nanoprecipitates is governed by the laser intensity of WLSP and processing temperature. Similarly, the uniformly distributed high density of dislocations is formed by WLSP and is influenced by the interaction of the gliding

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50 nm

50 nm Uniformly distributed dislocations

200 nm

Fig. 4.28 TEM images showing microstructures after WLSP: highly dense spherical nanoprecipitates (pointed out by red arrows) are generated in AA6061, and highly dense and uniformly distributed dislocations are entangled with nanoprecipitates in AA6061. (Liao, Y., Ye, C., Gao, H., Kim, B.-J., Suslov, S., Stach, E.A. et al., 2011. J. Appl. Phys. 110, 023518. With kind permission of AIP Publishing.)

dislocations with the nanoprecipitates (Fig. 4.29). The laser power and processing temperature effects on the mechanical properties expressed in terms of surface hardness are illustrated in Fig. 4.30. At this stage, it would be important to provide some expressions for the dislocation dynamics at the nano-micro level. Dislocation loops with arbitrary shapes interact with themselves, each other, and other crystal defects such as point (or cluster) defects, microcracks, and microvoids. The velocity and motion of a dislocation node can be obtained by solving a Newtonian equation consisting of an inertia term, a drag term, and a driving force vector presented in Eqs. (4.3) and (4.4) along with Liao et al. (2011).

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50 nm

50 nm

(a)

(b)

400 nm

(c) Fig. 4.29 TEM images showing microstructures after LSP at room temperature. (a) Nanoprecipitates are hardly observed in AA6061. (b) and (c) Dislocation pileups (pointed out by red arrows) are formed in AA6061. (Liao, Y., Ye, C., Gao, H., Kim, B.-J., Suslov, S., Stach, E.A. et al., 2011. J. Appl. Phys. 110, 023518. With kind permission of AIP Publishing.)

  1 dW mv_ 1 + vi ¼ Fi mi ¼ MiðT , P Þ vi dv ! N 1 X self σ Sj:j + 1 + σ other  bi  vi + Fi Fi ¼ 1

(4.3) (4.4)

j¼i

where mi is the effective dislocation segment mass density, Mi is the dislocation mobility determined by temperature T and pressure P, W is the total

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140

VHN (kg/mm2)

130 120 110 100 90

LSP at room temperature WLSP at 90°C

80

WLSP at 160°C 70

0.8

1.6

2.4

Laser power intensity (GW/cm2) Fig. 4.30 The effects of processing temperature and laser power intensity on surface hardness. (Y. Liao, Ch. Ye, H. Gao, B. -J. Kim, S. Suslov, E. A. Stach, and G. J. Cheng, J. App. Phys., 110, 023518 (2011). With kind permission of AIP Publishing.)

energy per unit length of moving dislocation, vi is the dislocation velocity, b is the Burgers vector, and Fi is the Petch-Kohler force on a dislocation node i, which is governed by three components: the stress from all dislocation P 1 S self other reploops and curves N j¼1 σ j:j + 1 , the self-force Fi and the term σ resenting the combined effects of the internal lattice friction, external applied stresses, and stresses due to other crystal defects such as the stacking-fault tetrahedral and Frank sessile loops. The macroscopic plastic strain rate ε_ p and the plastic strain Wp are related and determined by the motion of each dislocation segment according to:

ε_ p ¼

N X li vgi i¼1

Wp ¼

2V

ðni ⨂bi + bi ⨂ni Þ

N X li vgi i¼1

2V

ðni ⨂b1  bi ⨂ni Þ

(4.5)

(4.6)

where li is the dislocation segment length, vgi is the glide velocity, ni is a unit normal to the slip plane, V is the volume of the representative volume element, and N is the total number of dislocation segments, respectively.

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σother of Eq. (4.4) represents the contribution of all the Frank sessile loops in the multiscale discrete dislocation dynamic (MDDD) stimulation. By integrating the Peach-Kohler equation for self-stress of any curved closed dislocation loop, an expression of the stress field of one such Frank sessile loop can be written as x σ xz ¼ C 0 ½D1 E ðkÞ + D2 K ðkÞ ρ σ zz ¼ C 0 ½D3 E ðkÞ + D4 K ðkÞ xy fC ½D5 EðkÞ + D6 K ðkÞ + C 0 ½D7 E ðkÞ + D8 K ðkÞg ρ2 ρxx ¼ C ½D9 EðkÞ + D10 K ðkÞ + C 0 ½D11 EðkÞ + D12 K ðkÞ

σ xy ¼

(4.7)

The coefficients D1–D16 are used to simplify the stress expressions and are functions of spatial coordinates. They include the loop radius R, the cylindrical coordinates ρ and z, and the position vector of a field point  1=2 r ¼ ρ 2 + z2 (4.8) Expression for D1–D16 can be found in the works of Khraishi. K(k) and E(k) are the complete elliptic integrals of the first and second kind from the PK equation, and k is the modulus of these integrals. The parameters C, C0 , a, b, and k are defined in Eq. (4.7) as Gbz Gbz C¼ , C0 ¼  , a ¼ r 2 + R2 , b ¼ 2ρR, ð Þ z 2π 1  v   2b 1=2 k¼ a+b

(4.9)

Here bz is the elastic constant, G is the shear modulus, and ν is the Poisson ratio, respectively. Note that the symbol ⨂ in Eqs. (4.5) and (4.6) is the Kronecker product which in mathematics is an operation on two matrices of arbitrary size resulting in a block matrix. For details on the use of WLSP and of the MDDD simulation for AA6061, the reader is referred to the work of Liao et al. (2011). Summarizing the dislocation pinning in AA6061 in the light of WLSP and MDDD simulation, the following concepts emerge. (1) Due to the interaction between gliding dislocations and the dense nanoprecipitates, high density of dislocations are formed after WLSP which are uniformly distributed. (2) The pinning of dislocations by the high volume density of the

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Fig. 4.31 TEM image showing dislocations bowing around precipitates in the parent material, but no Orowan loop formation is observed. (Turnage, S.A., Hirth, J.P., Rajagopalan, M., Whittington, W.R., Tschopp, M.A., Peralta, P. et al., 2018. Mater. Sci. Eng. A 724, 609. With kind permission of Elsevier.)

nanoprecipitates depends on temperature and the laser intensity of WLSP. (3) The size, volume density, and the space distribution of the nanoprecipitates influence the effectiveness of the pinning of the dislocations. It would be of interest to indicate (Fig. 4.31) that precipitates in the parent material strengthen the alloy and dislocations bowing around occur when cutting through the precipitate is hindered. 4.2.2.2 The strain effect in 6061 Dislocation cells with different dislocation densities and cell sizes are formed when at a constant strain (0.25) deformation is performed at different strain rates. At Fig. 4.32a deformed at 103 s1, loosely tangled dislocation cells with a large number of dislocation loops within the cell interiors and cell walls are seen. The deformation was impact-loaded and the formation of the dislocation cells is consistent with the high stacking fault energy of aluminum which facilitates cell formation by cross-slip. Specimens deformed at 2  103 s1, seen in Fig. 4.32b, also consist of dislocation cells, but the dislocation cells are better defined. Numerous individual dislocations and dislocation loops can be resolved within the cell walls. Some isolated dislocation loops are also observed within the cell interiors. When the strain rate is 3  103 s1 the specimens exhibit very well-defined cell walls with a high density of dislocations within the walls. Isolated dislocation loops are still observed within cell interiors. Thus, it was observed that substructure

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Fig. 4.32 Dislocation substructure of specimens deformed at strain rates of (a) 103 s1, (b) 2  103 s1, (c) 3  103 S1, and (d) 4  103 s1 at a true strain of 0.25. (Lee, W.-S., Shyu, J.-C., Chiou, S.-T., 2000. Scr. Mater. 42, 51. With kind permission of Elsevier.)

evolution consists of dislocation cells and dislocation loops. It could be noted that the dislocation density and cell size are directly affected by the strain rate level. The microstructural changes considered above are influenced by the strain and the strain rate affecting directly the dislocation level and, consequently, the plastic flow stress behavior.

4.3 In stainless steel 304L 4.3.1 Introduction The mechanical properties of the austenitic SS 304L can be enhanced by deformation. In the wake of deformation, the dislocation density of the base metal increases. The strengthening effect of materials processed by large strain deformation is a consequence of grain boundary and dislocation strengthening. As known and discussed earlier, the strengthening by dislocations depends on the square root of the dislocation density given as

Dislocations in AM and traditional manufacturing

σ disl ¼ αGbρ0:5

205

(4.10)

where α, G, and b are constants representing the shear modulus and Burgers vector and ρ is the dislocation density. Grain size and dislocation strengthenings can be combined to the relation of σ 0:2 ¼ σ 0 + Kε D0:5 + αGBρ0:5

(4.11)

Here σ 0 is the strength of the dislocation-free single crystal, whereas KεD0.5 stands for the grain boundary strengthening αGB with D being the grain size and Kε a constant. It is usually assumed that grain boundary and dislocation strengthening contribute independently to the strength. Here we consider only the dislocation patterns that evolve in SS 304L.

4.3.2 In AM 304L stainless steel Measured dislocation density in 304L SS as 7.45  1012 m2 and 4.31  1012 m2 for the area is shown in Fig. 4.33a and b, respectively. The alloy was produced by the laser engineered net shape (LENS) and in the figure TEM micrographs are shown. Dislocations can be seen accumulating around the particles in the microstructure. They also contribute to the strengthening during plastic deformation in addition to the dislocations. STEM micrograph of 304L obtained by electro optical system (EOS) is shown in Fig. 4.34. High dislocation density of 2.72  1014 m2 was

Fig. 4.33 (a) Bright-field and (b) dark-field STEM microstructures of 304L stainless steel LENS build showing fine dispersion of submicron steel particles, and (c) a typical particle. The arrows in (a) show examples of the fine particles throughout the microstructure [shown larger in (c)]. The arrow in (b) shows a large impurity, which is much less common though still present in the microstructure. (Morrow, B.M., Lienert, T.J., Knapp, C.M., Sutton, J.O., Brand, M.J., Pacheco, R.M., et al., 2018. Metall. Mater. Trans. A 49A, 3637. With kind permission of Springer Nature.)

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Fig. 4.34 (a) Bright-field and (b) dark-field STEM microstructures of 304L stainless steel EOS build showing fine dispersion of sub-micron-sized particles. (Morrow, B.M., Lienert, T.J., Knapp, C.M., Sutton, J.O., Brand, M.J., Pacheco, R.M., et al., 2018. Metall. Mater. Trans. A 49A, 3637. With kind permission of Springer Nature.)

observed for the area shown in Fig. 4.34a. The difference in the dislocation content (higher dislocation density) may be the result of the faster cooling rate and the higher residual stresses of the EOS technique than in the LENS built microstructure. The particles observed in the microstructure are generally round in shape. A comparison between the micrographs of wrought 304L SS not fabricated by AM is shown in Fig. 4.35 for comparison purposes. A high dislocation density of 1.84  1014 m2 similar to that observed in

Fig. 4.35 Bright-field STEM micrograph of 304L stainless steel produced through a traditional wrought processing technique. No fine dispersion of particles is evident. (Morrow, B.M., Lienert, T.J., Knapp, C.M., Sutton, J.O., Brand, M.J., Pacheco, R.M., et al., 2018. Metall. Mater. Trans. A 49A, 3637. With kind permission of Springer Nature.)

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Table 4.1 Measured dislocation density for the microstructural areas shown in micrographs of AM and wrought 304L and 316L stainless steels shown in Figs. 4.33–4.35. Material

Dislocation density (m22)

316L 316L 304L 304L 304L

Too low to quantify 2.77  1014 m2 4.31 to 7.45  1012 m2 2.72  1014 m2 1.84  1014 m2

LENS HT LENS AΒ LENS AB EOS AB Wrought

No figures of 316L are shown. Morrow, B.M., Lienert, T.J., Knapp, C.M., Sutton, J.O., Brand, M.J., Pacheco, R.M., et al., 2018. Metall. Mater. Trans. A 49A, 3637. With kind permission of Springer Nature.

the structure by EOS was obtained. But it is not a consequence of the residual stress of EOS-built material but rather the result of the extensive cold work associated with the wrought forged process. Table 4.1 list the measured dislocation densities in 304L SS alloys. Alloys of 3.16L are also included.

4.3.3 In conventionally fabricated 304L stainless steel In this section, dislocation patterns that evolved in the 304L base metal stainless steel (SS) during cyclic deformation are presented as observed by TEM. Fig. 4.36 illustrates the dislocation structure of the undeformed base 304L SS. In this condition, 3304L is characterized by planar dislocation structures: discrete dislocation lines in Fig. 4.36a, pileups in Fig. 4.36b, tangles in Fig. 4.36c, and stacking faults in Fig. 4.36c and d. The dislocation density is extremely low in this initial condition. The dislocation structure after strain cycling with Nint ¼ 15 is shown in Fig. 4.37 (Nint indicates the number of cycles at interruption in the cycles on deformation). The microstructures after 100 cycles, namely, Nint ¼ 100, are seen in Fig. 4.38, while after Nint ¼ 1000 they are illustrated in Fig. 4.39. In Fig. 4.40, the dislocation structure of Nint ¼ 2526 cycles is presented. In the structures 4.36–4.40 shown after 15 cycles, the primary hardening stage ends and the dislocations are still in the form of planar structures but their density increase significantly. The dislocation interactions are complex. The planar structures are tangled as seen in Fig. 4.37a and forming incipient dislocation walls and dipoles are also formed. Zigzag dislocations are observed in low density regions, which indicates that cyclic plastic deformation is significantly influenced by solute atoms (atmosphere) such as C or N. They restrict the motion of the

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Fig. 4.36 Microstructures in the undeformed 304L base metal: (a) discrete dislocation lines in grain, (b) dislocation pileup, (c) light dislocation tangles and stacking faults nearby the grain boundary, (d) long-range stacking faults. (Wang, H., Jing, H., Zhao, L., Han, Y., Lv, X., Xu, L., 2017. Mater. Sci. Eng. A 690, 16. With kind permission of Elsevier.)

dislocations. Crossed stacking faults are often observed and the amounts of the stacking faults increase as seen in Fig. 4.37b. In some grains, deformation twins are formed. At the deformation twin boundaries, dislocations and stacking faults are piled and tangled (Fig. 4.37c). After 100 cycles, which is near the end of the softening stage, dislocations wall/channel structures are observed as shown in Fig. 4.38a. In the wall/channel structure, the dislocation walls are formed by high-density dislocation tangles and channel with low density of dislocations. In the last hardening stage, interaction between the dislocations, stacking faults, and deformation twins induces shear bands gradual formation intermixed with dislocations, stacking faults, and deformation

Fig. 4.37 Microstructures of 304L base metal during strain cycling (Nint ¼ 15c): (a) incipient dislocation walls, (b) stacking faults, (c) deformation twins. (Wang, H., Jing, H., Zhao, L., Han, Y., Lv, X., Xu, L., 2017. Mater. Sci. Eng. A 690, 16. With kind permission of Elsevier.)

Fig. 4.38 Microstructure of 304L base metal during strain cycling (Nint ¼ 100c): (a) dislocation wall/channel structures, (b) shear bands with interior dislocation veins and carpet-like structure dislocation tangles. (Wang, H., Jing, H., Zhao, L., Han, Y., Lv, X., Xu, L., 2017. Mater. Sci. Eng. A 690, 16. With kind permission of Elsevier.)

Fig. 4.39 Microstructures of 304L base metal during strain cycling (Nint ¼ 1000c): (a) dislocation cells, (b) α0 –martensite in the shear bands. (Wang, H., Jing, H., Zhao, L., Han, Y., Lv, X., Xu, L., 2017. Mater. Sci. Eng. A 690, 16. With kind permission of Elsevier.)

Fig. 4.40 Microstructure of 304L base metal during strain cycling (Nint ¼ 2526c): (a) dislocation lines in labyrinth structure, (b) dislocation veins in cells, (c) α0 –martensite in the shear bands. (Wang, H., Jing, H., Zhao, L., Han, Y., Lv, X., Xu, L., 2017. Mater. Sci. Eng. A 690, 16. With kind permission of Elsevier.)

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twins as shown in Fig. 4.38b. Inside the shear bands, dislocation vein and highdensity carpet-like structure dislocation are observed. During 1000 cycles (it is the second stage hardening), dislocation cells are observed—some of them not well-developed—as seen in Fig. 4.39a. In some grains, the amount of shear bands and interior dislocations density greatly increase and also α0 —martensite is nucleated inside the shear bands (deformation induced martensite transformation) as seen in Fig. 4.39b. At the 2526 cycle (the end of fatigue life) labyrinth structure, dislocation cells and dislocation lines and veins still exist in their interior zones as shown in Figs. 4.40a and b. Further, in the shear bands, the amount and size of martensite increase with increasing cycle number (Fig. 4.40c).

4.4 In alumina (Al2O3) Limited information exists on additive manufactured pure alumina. Despite that, AM offers many advantages over traditional processing. Among them, fabricating complex parts of intricate shape and size with the same cost (or even reduced cost), no pure dislocation structures and their contribution to the mechanical properties are covered in research data in the AM field. The reason might be that alumina is used on a large scale to strengthen metallic matrixes, as a component of other ceramics to improve properties (e.g., zirconia, SiC, Yag, etc.) and probably also because flaws and large porosity are very often obtained as a result of the AM process. It is well known that the properties, in particular the mechanical of ceramics, are very sensitive to porosity resulting in reduced performance, namely, lower resistance to stress. Producing monolithic ceramics, among them alumina enabling them to match the physical properties (in particular, the mechanical) of the traditionally manufactured counterparts, is still a challenge. Research is being continuously done to improve methods for the fabrication of dense ceramics, in particular those of intricate shapes by the AM technique which is an important task facing AM technology. Of the many AM techniques used to fabricate components, extrusion-based and lithography-based AM processes are promising, because they are capable of producing >95% dense ceramic parts. Because of the lack of dislocation images in monolithic alumina, we shall proceed forward right away to the next section of conventionally fabricated alumina.

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4.4.1 In conventionally fabricated alumina Relatively little TEM information has been performed on the hard ceramic materials compared with metals and alloys. The main reason is the difficulty in preparing thin enough sections to be electron transparent. Whichever technique is used, care should be taken not to alter the dislocation configuration during the annealing out surface damage. Early work on the direct observation of dislocations in alumina is illustrated in Fig. 4.41. Dislocation in unannealed single-crystal alumina after bending can be seen at low magnification in Fig. 4.41a. In Fig. 4.41b, alternating black and white contrast at

Fig. 4.41 Dislocations in alumina: (a) General low magnification field; (b) alternating black and white contrast in A, zigzag contrast at B and a node in C; and (c) alternating contrast often a bifurcated terminal region. (Evans, P.E., Hardiman, B.P., 1965. Nature 4980, 182. With kind permission of Springer Nature.)

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A, zigzag contrast at B and possible node at C are seen, and in Fig. 4.41c again alternating contrast is seen often with a bifurcated terminal region. This contrast is not well understood and cannot be unequivocally associated with dislocation dissociation to partials. More illustrative dislocations in the alumina phase can be seen in a unidirectionally solidified eutectic Al2O3-YAG composite. In the microstructure, single-crystal Al2O3 and single-crystal YAG are continuously connected, finely entangled without grain boundaries. Only the dislocation structure in the alumina phase is illustrated in Figs. 4.42 and 4.43 of the plastically deformed specimens at high temperatures, in both tensile and compressive tests, respectively. The plastic deformation in the single-crystal composite occurred by dislocation motion. More dislocations of linear nature were observed in the Al2O3 single crystal than in the YAG single crystal of the composite specimen. The composite material exhibits excellent creep resistance. The dislocation structure in crept specimen loaded

Fig. 4.42 TEM image showing the dislocation structures of the Al2O3 phase. Specimen plastically deformed after tensile test at 1700°C of the Al2O3-YAG single-crystal composite. (Waku, Y., Sakuma, T., 2000. J. Eur. Ceram. Soc. 20, 1453. With kind permission of Elsevier.)

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Fig. 4.43 TEM image showing the dislocation structure of the Al2O3 phase. The illustration is of the sintered composite specimen compressively crept at 1600°C at a strain rate of 105/s. (Waku, Y., Sakuma, T., 2000. J. Eur. Ceram. Soc. 20, 1453. With kind permission of Elsevier.)

compressively at the high temperature of 1600°C is seen in Fig. 4.43. This single-crystal Al2O3-YAG composite has superior high-temperature strength in tensile test at the above 1650°C (flexural test shows no temperature dependence in the range from room temperature up to 1800°C) and the marked plastic deformation occurs by dislocation motion. The compressive flow stress of the single-crystal Al2O3-YAG composite is about 13 times higher than that of a sintered material with the same composition. Only the portion related to the Al2O3 dislocation structure has been shown of this Al2O3-YAG composite material. It is important to show two additional aspects of the dislocation structure in alumina, namely, (a) dipole structure and loop formation, and (b) the dislocation configuration in low-angle boundaries which is an array of dislocations. (a) In many metals, alloys, and ceramic materials, dipoles are observed after deformation on a single slip system. This occurs when two dislocations

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215

of opposite sign glide on parallel slip planes are trapped by their mutual stress fields forming the dipole. Edge dislocation dipoles formed in sapphire during basal glide break up into loop by a series of processes involving self-climb. Easy glide deformation in HCP occurs by basal plane. Sapphire is an α-alumina single crystal and it deforms on basal plane. High-temperature deformation is required in sapphire to enable diffusion to take place and promote breakup of dislocation dipoles into loops. A dislocation substructure after deformation on the (0001) [1120] basal glide system at 1400°C is illustrated in Fig. 4.44. The substructure consists of glide dislocations predominantly in edge orientation, edge dislocation dipoles, and dislocation loops. Screw dislocations annihilate by cross-slip because they are assumed

Fig. 4.44 Typical dislocation structure in sapphire deformed by basal slip at 1400°C. Dislocation dipoles and loops formed from dipoles are a major component of the dislocation density (0001) foil, 650 kV. (Phillips, D.S., Pletka, B.J., Heuer, A.H., Mitchell, T.E., 1982. Acta Metall. 30, 498. With kind permission of Elsevier.)

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to have a higher velocity than edge dislocations. The continuous breakup of the edge dipoles into loops by self-climbing during deformation and the competition between dipole formation and annihilation determine the work hardening and recovery of sapphire. Doping by Ti4+ affects the diffusion kinetics in such a way that strings of perfect loops are found as observed in Fig. 4.45. The theoretical evaluation of the breakup of the dipoles and loop formation, the kinetics, and the energetics are widely discussed in the work of Phillips et al. (1982). (b) Low-angle grain boundaries or subgrain boundaries (also small angle boundaries) are composed of an array of dislocations. Grain boundaries falling in this category have a misorientation . (Tochigi, E., Nakamura, A., Shibata, N., Ikuhara, Y., 2018. Crystals 8, 133; Open access.)

{1120}/[0001] 2° low-angle tilt boundary is shown. Pair contrasts are periodically arranged suggesting that each dislocation is dissociated into two partial dislocations with Burgers vectors given as b ¼ 1/3 because the translation vector of 1/3 is perpendicular to the {1120} plane. The relationship between the interval of perfect dislocations d, the Burgers vector b, and the misorientation angle of a low-angle grain boundary θ is given by Frank’s equation d ¼ |b|=θ

(4.12)

HRTEM image of a dislocation pair in the grain boundary is shown in Fig. 4.46b. These partial dislocations separated along the {1120} plane and a stacking fault is formed between the partial dislocations. The plane of the stacking fault does not coincide with the slip plane of the perfect dislocation. The partial dislocations were separated by a self-climb mechanism. The partials are formed by the reaction: 1=3 < 1120 >! 1=3 < 1010 > + 1=3 < 0110 >

(4.13)

In Fig. 4.47a, TEM image of {1120}/ low-angle 2° tilt grain boundary is seen. The dislocations are arranged periodically along the grain boundary. The dislocation structure is divided into two groups, pairs of dislocations, and groups of odd-numbered dislocations. The large Burgers circuit in Fig. 4.47c shows that the dislocation pair has the Burgers vector of 1/3 in total. The small Burgers circuits show that each partial

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Fig. 4.47 (a) TEM image of the {1120}/ 2° low-angle tilt grain boundary. The grain boundary consists of partial-dislocation pairs and groups of 5 and 13 partial dislocations. (b) Dark-field TEM image taken at the same region in (a) using g ¼ 3030, where the grain boundary plane is inclined by about 30°. Open and filled triangles indicate 1/3 and 1/3 < 0110> partial dislocations, respectively. (c) HRTEM image of a partial-dislocation pair viewed along the [1100] zone axis. From the Burgers circuits, it is found that the dislocation pair has the Burgers vector of 1/3 < 1120 > and the partial dislocations have 1/3 < 1010> and 1/3 < 0110>. (Tochigi, E., Nakamura, A., Shibata, N., Ikuhara, Y., 2018. Crystals 8, 133; Open access.)

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dislocation has an edge component of 1/6 . This component corresponds to the {1100} projection of the vectors of 1/3 and 1/3 . Therefore, it is considered that the observed structure corresponds to the dissociation of the 1/3 edge dislocation into the 1/3 and 1/3 mixed partial dislocations according to the reaction of Eq. (4.13). In Fig. 4.47b, the partial dislocations with b1 ¼ 1/3 and with b2 ¼ 1/3 have strong and weak contrast, respectively. The edge components of the partial dislocations with b1 and b2 are both 1/6 , whereas their screw components are 1/6 and 1/6 , respectively. It has been suggested that the odd-numbered dislocation structures are generated by the twist component of the grain boundary. TEM images of the {1100}/ 2° tilt boundary are illustrated in Fig. 4.48. Dislocation triplets are arranged along the grain boundary, suggesting that the dislocations dissociated into three partials with {1100} stacking faults. HRTEM image of triplet is seen in Fig. 4.48b. The Burgers circuits show that the dislocation triplet has the Burgers vector of < 1100 > in total. Dislocations slip is prism-plane and the dissociation reaction is written as < 1100 >! 1=3 < 1100 > + 1=3 < 1100 > + 1=3 < 1100 > Illustrations of dislocation structures in other low-angle grain boundaries listed in Table 4.2 are not shown here. In the work of Tochigi et al. (2018), details on structure of the low-angle grain boundaries (mainly tilt boundaries but one twist boundary is also included), the dissociation of perfect dislocations into partials, and the stacking sequence of the stacking faults formed between the partials are thoroughly discussed and are not further discussed here, except for the listing of the dislocation orientations observed in the low-angle boundaries of the bicrystals: (1) {1120}/[0001] tilt grain boundary The 1/3 perfect edge dislocation dissociates into 1/3 < 1010 > and 1/3 partial dislocations with the {1120} stacking fault. (2) {1120}/ tilt grain boundary The 1/3 perfect edge dislocation dissociates into 1/3 < 1010 > and 1/3 partial dislocations with the {1120} stacking fault. (3) {1100}/ tilt grain boundary The perfect edge dislocation dissociates into 1/3 partial dislocations with the {1100} stacking faults of I2 and V.

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Fig. 4.48 (a) TEM image of the {1100}/ 2 low-angle tilt grain boundary. The grain boundary consists of dislocation triplets. (b) HRTEM image showing one of the dislocation triplets. The Burgers circuit indicates the Burgers vector of < 0110 >, suggesting that the < 1100 > dislocation is dissociated into 1/3 partial dislocations with two stacking faults in between. (c) ABF STEM image of the left stacking fault in (b). The atomic structure model overlapped with the image is shown at the right panel. The stacking sequence is …ABC/C/AB…: I2. (d) The right stacking fault in (b). The stacking sequence is …ABC//BCAB. (Tochigi, E., Nakamura, A., Shibata, N., Ikuhara, Y., 2018. Crystals 8, 133; Open access.)

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(4) /[1100] tilt grain boundary The misorientation of the grain boundary is accommodated by the groups of the 1/3[0111], 1/3[1011], and 1/3[1101] perfect mixed dislocations. Each perfect dislocation dissociates into 1/18 and 1/18 partial dislocations with the (0001) stacking fault. (5) {1104}/ tilt grain boundary The misorientation of the grain boundary is accommodated by 1/3 perfect mixed dislocations and 1/3 perfect edge dislocations. The 1/3 dislocations are not dissociated, whereas 1/3 edge dislocations dissociate into 1/18 and 1/18 with the (0001) stacking fault. (6) (0001)/[0001] twist grain boundary In the hexagonal network the 1/3 prefect screw dislocations are formed. The 1/3 screw dislocation is not dissociated into partial dislocations. Note: According to the stacking disorder, stacking faults are often called: interstitial fault type-I (I1), interstitial fault type-II (I2), and vacancy fault (V), respectively.

References Liao, Y., Ye, C., Gao, H., Kim, B.-J., Suslov, S., Stach, E.A., Cheng, G.J., 2011. J. Appl. Phys. 110, 023518. Pelleg, J., 2012. Mechanical Properties of Materials. Springer. Pelleg, J., 2014. Mechanical Properties of Ceramics. Springer. Phillips, D.S., Pletka, B.J., Heuer, A.H., Mitchell, T.E., 1982. Acta Metall. 30, 498. Tochigi, E., Nakamura, A., Shibata, N., Ikuhara, Y., 2018. Crystals 8, 133.

Further reading Castany, P., Pettinari-Sturmel, F., Crestou, J., Douin, J., Coujou, A., 2007. Acta Mater. 55, 6284. Evans, P.E., Hardiman, B.P., 1965. Nature 4980, 182. Khraishi, T.A., Hirth, J.P., Zbib, H.M., de La Rubia, T.D., 2000. Philos. Mag. Lett. 80 (2), 95. Khraishi, T.A., Zbib, H.M., de la Rubia, T.D., Victoria, M., 2001. Philos. Mag. Lett. 81, 583. Lee, W.-S., Shyu, J.-C., Chiou, S.-T., 2000. Scr. Mater. 42, 51. Liu, C.Y., Qu, B., Xue, P., Ma, Z.Y., Luo, K., Ma, M.Z., Liu, R.P., 2018. J. Mater. Sci. Technol. 34, 112. Morrow, B.M., Lienert, T.J., Knapp, C.M., Sutton, J.O., Brand, M.J., Pacheco, R.M., Livescu, V., Carpenter, J.S., Gray III, G.T., 2018. Metall. Mater. Trans. A 49A, 3637. Murr, L.E., Esquivel, E.V., Quinones, S.A., Gaytan, S.M., Lopez, M.I., Martinez, E.Y., Medina, F., Hernandez, D.H., Martinez, E., Martinez, J.L., Stafford, S.W.,

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Brown, D.K., Hoppe, T., Meyers, W., Lindhe, U., Wicker, R.B., 2009a. Mater. Charact. 60, 96. Murr, L.E., Quinones, S.A., Gaytan, S.M., Lopez, M.I., Rodela, A., Martinez, E.Y., Hernandez, D.H., Martinez, E., Medina, F., Wicker, R.B., 2009b. J. Mech. Behav. Biomed. Mater. 2, 20. Murr, L.E., Gaytan, S.M., Medina, F., Lopez, H., Martinez, E., Machado, B.I., Hernandez, D.H., Martinez, L., Lopez, M.I., Wicker, R.B., Bracke, J., 2010. Phil. Trans. R. Soc. A 368, 1999. Murr, L.E., Gaytan, S.M., Ramirez, D.A., Martinez, E., Hernandez, J., Amato, K.N., Shindo, P.W., Medina, F.R., Wicker, R.B., 2012. J. Mater. Sci. Technol. 28(1). Summers, P.T., Chen, Y., Rippe, C.M., Allen, B., Mouritz, A.P., Case, S.W., Lattimer, B.Y., 2015. Fire Sci. Rev. 4, 3. Turnage, S.A., Hirth, J.P., Rajagopalan, M., Whittington, W.R., Tschopp, M.A., Peralta, P., Solanki, K.N., 2018. Mater. Sci. Eng. A 724, 609. Waku, Y., Sakuma, T., 2000. J. Eur. Ceram. Soc. 20, 1453. Wang, H., Jing, H., Zhao, L., Han, Y., Lv, X., Xu, L., 2017. Mater. Sci. Eng. A 690, 16. Woo, W., Unga´r, T., Feng, Z., Kenik, E., Clausen, B., 2010. Metall. Mater. Trans. A 4, 1210. Yan, C., Hao, L., Hussein, A., Young, P., 2015. J. Mech. Behav. Med. Mater. 51, 61.

CHAPTER FIVE

Deformation in AM and traditional manufacturing: A comparison

5.1 Introduction Plastic deformation is the permanent distortion that occurs when a material is subjected to stress beyond its yield strength. The stress can be tensile, compressive, bending, or torsion and when the yield stress is exceeded the sample either elongates, get compressed (in ductile material to a barrel shape known as barreling), buckle, bend, or twist according to the type of deformation, respectively. The deformation is temperature dependent, but at a constant temperature it depends on the mobility of the structural defects such as dislocations; but the mobility also depends on the grain boundaries (in polycrystals), stacking faults, twins, and point defects. All these factors determine the extent of the plastic deformation, which is generally described as strain. Clearly, since temperature is an important factor, the rate of atomic diffusion is a limiting factor on the mobility of the defects. Plastic deformation is characterized by irreversible strain after the removal of the applied force, contrary to elastic deformation (also known as temporary deformation), which is recoverable after removal of the applied force. Under the applied force, plastic deformation is characterized by work hardening (stain hardening), which is followed by a necking region (usually after the UTS is reached) until fracture sets in as the final stage of the deformation. During the strain-hardening stage, the material becomes stronger as a consequence of dislocation motion. During necking, the cross section of the specimen is reduced and the reduced area of the specimen can no longer withstand the applied load (at this stage, the stress is at its maximum) and the strain rapidly increases to fracture.

Additive and Traditionally Manufactured Components https://doi.org/10.1016/B978-0-12-821918-8.00005-X

© 2020 Elsevier Inc. All rights reserved.

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5.1.1 Deformation in AM Ti-6Al-4V It has been mentioned earlier when dealing with Ti-6Al-4V that the wide interest in this alloy is related to its use in aerospace, nuclear industry, but not at least in the biomedical industry as implants. This wide use of Ti-6Al-4V is due to its excellent properties such as high specific strength, low density, and exceptionally good corrosion resistance. The biomedical use of Ti-6Al-4V is possible because its biocompatibility. Recent interest has shifted from conventional manufacturing to any of the many possible additive fabrication methods for producing various materials among them Ti alloys. Choosing the most suitable and flexible AM method is essential to get the best quality product. Selective laser melting (SLM) is one of them, which has gained wide interest both in industrial application and research due to its high level of flexibility. Stress-strain relation of Ti-6Al-4V alloy—prepared by SLM obtained—by compression is illustrated in Fig. 5.1 at two strain rates. Strain softening of the flow stress of the samples tested at 950°C sets in immediately after yielding at a low strain below
0.05. As seen in the figure, there is an increase in flow stress with increase in the stain rate. This observation is probably associated with the lack of enough time for dynamic recrystallization softening to

Fig. 5.1 Compressive stress-strain curves of Ti-6Al-4V samples produced by SLM process at different strain rates. (Zhang, Q., Liang, Z.-L., Cao, M., Liu, Z.-F., Zhang, A.-F., Lu, B.-H., 2017. Trans. Nonferrous Metals Soc. China 27, 1036. With kind permission of Elsevier.)

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(a)

(b)

Porosity Porosity

20 mm

20 mm (d)

(c)

1 mm

0.2 mm

Fig. 5.2 Micrographs of as-built Ti-6Al-4V specimens: (a) horizontal view; (b) vertical view; and (c and d) high magnification of horizontal section. (Zhang, Q., Liang, Z.-L., Cao, M., Liu, Z.-F., Zhang, A.-F., Lu, B.-H., 2017. Trans. Nonferrous Metals Soc. China 27, 1036. With kind permission of Elsevier.)

occur. The micrograph of the as-built Ti-6Al-4V alloy is shown in Fig. 5.2, while in the SEM microstructures of Fig. 5.3 forged and water cooled followed by the SLM process are indicated. The letters A–D in the figure indicate the process parameters such as strain and strain rates as listed in Table 5.1. Since the mechanical properties of material depend on the porosity content of the samples, evaluation of its amount is desired. This is shown in Fig. 5.4. Further, the microhardness of these samples is illustrated in Fig. 5.5. Built Ti-6Al-4V by additive manufacturing using electron beam melting (EBM) technique is another method to produce components and it has a great potential, among other, for implant production. The yield strength, the ultimate tensile stress, and the elongation of the specimens fabricated by EBM technique are listed in Table 5.2. The variation of the true stress with true strain is seen in Fig. 5.6. The work-hardening exponent n is shown for all specimens.

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(b)

(a)

10 mm

10 mm (c)

(d)

10 mm

10 mm (e)

(f)

10 mm

10 mm

Fig. 5.3 Micrographs of Ti-6Al-4V specimens following SLM processing, subsequent forged and water quenching: (a) horizontal view of sample B; (b) vertical view of sample B; (c) horizontal view of sample C; (d) vertical view of sample C; (e) horizontal view of sample D; and (f) vertical view of sample D. (Zhang, Q., Liang, Z.-L., Cao, M., Liu, Z.-F., Zhang, A.-F., Lu, B.-H., 2017. Trans. Nonferrous Metals Soc. China 27, 1036. With kind permission of Elsevier.) Table 5.1 An overview of process parameters for different samples. Sample no.

_ 21 Strain rate, ε=s

Strain, ε

A B C D

0 0.1 0.01 0.1

0 0.7 0.7 0.35

Zhang, Q., Liang, Z.-L., Cao, M., Liu, Z.-F., Zhang, A.-F., Lu, B.-H., 2017. Trans. Nonferrous Metals Soc. China 27, 1036. With kind permission of Elsevier.

2.0

Proportion of porosity

Horizontal section Vertical section 1.5 1.30 1.14 0.93

1.0 0.35

0.74 0.5

0.70

0.33

0.79

Sample B

Sample C

0 Sample A

Sample D

Fig. 5.4 Proportion of porosity at horizontal section and vertical section. (Zhang, Q., Liang, Z.-L., Cao, M., Liu, Z.-F., Zhang, A.-F., Lu, B.-H., 2017. Trans. Nonferrous Metals Soc. China 27, 1036. With kind permission of Elsevier.)

460 Microhardness on horizontal section Microhardness on vertical section

434.9

Microhardness (HV)

420 405.4

403.0 380.0

374.4 374.3

380

363.4 353.9

340

300 Sample A

Sample B

Sample C

Sample D

Fig. 5.5 Microhardness on horizontal section and vertical section. (Zhang, Q., Liang, Z.-L., Cao, M., Liu, Z.-F., Zhang, A.-F., Lu, B.-H., 2017. Trans. Nonferrous Metals Soc. China 27, 1036. With kind permission of Elsevier.) Table 5.2 Yield strength, ultimate tensile strength and elongation of four EBM-built Ti-6Al-4V samples with different build heights. Specimen

Yield strength (0.2% offset) (MPa)

Ultimate tensile stress (MPa)

Elongation (%)

10 mm1 10 mm2 10 mm3 10 mm4

851.8  5.8 836.6  8.7 827.9  0.3 823.4  0.1

964.5  0.3 953.7  4.3 944.5  5.8 940.5  6.5

16.3  0.8 15.2  1.2 14.0  0.5 13.2  0.7

Tan, X., Kok, Y., Tan, Y.J., Descoins, M., Mangelinck, D., Tor, S.B. et al., 2015. Acta Mater. 97, 1. With kind permission of Elsevier.

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900 10mm-1

True Stress (MPa) in Log

890

n1=0.132

10mm-2 10mm-3

880

10mm-4

n2=0.135

870 n3=0.136

860 850

n4=0.141 840 830 820 0.008

0.0085

0.009

0.0095

0.01

0.0105

True Strain in Log Fig. 5.6 Variation of true stress-true stain curves for the plastic response of the four EBM-built samples, plotted on logarithmic axes. The measured power-law strain-hardening exponents are indicated for each sample. (Tan, X., Kok, Y., Tan, Y.J., Descoins, M., Mangelinck, D., Tor, S.B. et al., 2015. Acta Mater. 97, 1. With kind permission of Elsevier.)

Recall that σ ¼ Kεn

(5.1)

where Κ is the material strength coefficient and n is the strain-hardening (work-hardening) coefficient and its value is in the range 0–1. At n ¼ 0 the material is perfectly plastic, while at n ¼ 1 the material behaves completely elastic. The plot of Fig. 5.6 is on a logarithmic scale. Therefore, according to Eq. (5.1), the intercepts of the curves with the stress axis give the logarithm of the stress coefficient, while the slopes provide the stress exponents. The four tensile specimens listed in Table 5.2 were intended to investigate consistency of microstructure and mechanical properties of EBM-built parts. For this purpose, the test specimens (designated in Table 5.2) were cut one by one from a horizontal block and the designations are from bottom to top. Anisotropy is of interest in additive manufactured components. Experiments were designed to explore anisotropic tensile behavior in Ti-6Al-4V components. The specimens were fabricated with direct energy deposition additive manufacturing. Table 5.3 summarizes the mechanical properties and the designation of the specimens (appearing in the table) are indicated in Fig. 5.7.

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Table 5.3 Summary of mechanical properties. Transverse Transverse (upper tier) (lower tier) Longitudinal Baseplate

Tensile strength (MPa) 1041  12 1087  8 Yield strength (MPa) 945  13 970  17 Measured elongation (%) 18.7  1.7 17.6  0.7 Elongation per ASTM E8 (%) 14.5  1.2 13.6  0.5

1063  20 960  26 13.3  1.8 10.9  1.4

1053  7 975  6 21.2  0.6 17.3  0.5

In specimens L1–L8, the tensile axis of each specimen was aligned with the longitudinal direction. In specimens T1–T10, the tensile axis of each specimen was aligned with the transverse, or build, direction. Specimens T1–T5 were extracted from the upper half of the cruciform, which was not exposed to oxygen, and T6–T10 were extracted from the lower half, which was exposed to oxygen. The base specimens were extracted from the wrought baseplate. Reported values are average  standard deviation. Carroll, B.E., Palmer, T.A., Beese, A.M., 2015. Acta Mater. 87, 309. With kind permission of Elsevier.

Fabrication of the cruciform (see Fig. 5.7) was performed in an enclosed chamber purged with ultrahigh-purity argon to minimize oxygen contamination in the laser-deposited material. The mechanical properties of the component in longitudinal and transverse orientations with respect to the build direction were measured under uniaxial tension. Fig. 5.8 illustrates representative engineering stress-strain curves for the longitudinal, transverse, and baseplate specimens. Optical micrographs from near the bottom and near the top of the cruciform wall are illustrated in Fig. 5.9a and b, respectively. In both structures, fine Widmanst€atten structure is seen with lath width of the order of 1 μm. Only a small α0 (HCP martensite) appears to be present contrary to the previously reported as-fabricated AM Ti-6Al4V parts. The different observation in the current experiments is a consequence of the difference in the cooling rates from above the β transus. (Recall that Ti and Ti alloys can exist in two forms, an HCP alpha and a BCC beta. Ti transforms from its α to its β form at a temperature called the beta transus. The beta transus span the range from 700°C to 1050°C depending on the alloy composition.) The cooling rate was not sufficiently fast to form martensite. Consequently, the yield and tensile strengths are lower than observed in other works. However, even so, the strength values observed are well above the requirements for cast and wrought Ti-6Al-4V. As seen in Fig. 5.8, location-dependent and direction-dependent (also oxygen-dependent) mechanical properties are indicated. The anisotropy manifests itself in the elongation. Accordingly, the elongation along the transverse direction is notably higher than in the longitudinal direction (see Fig. 5.8). For getting isotropic properties, heat treatment is required. The isotropicity in Ti-6Al-4V manufactured by laser-powder-treatment is required. The isotropicity in Ti-6Al-4V manufactured by laser-powder deposition (LPD)

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Fig. 5.7 (a) Half of the as-fabricated AM cruciform on a 152 mm baseplate; walls are
100 mm tall. The vertical direction corresponds to the build, or transverse, direction. Brackets indicate the region of oxygen contamination. (b) Schematic showing how the tensile specimens were extracted from the component. The rectangular gauge sections of the tensile specimens were 4 mm  2 mm  9 mm. (Carroll, B.E., Palmer, T.A., Beese, A.M., 2015. Acta Mater. 87, 309. With kind permission of Elsevier.)

1200

Engineering stress (MPa)

1100 1000 900 800 700 600 500 0.00

Transverse (upper) Transverse (lower) Longitudinal Baseplate 0.05

0.15 0.10 Engineering strain

0.20

0.25

Fig. 5.8 Engineering stress-strain curves of representative specimens showing similar elastic responses, and highlighting the differences in strength and elongation values. (Carroll, B.E., Palmer, T.A., Beese, A.M., 2015. Acta Mater. 87, 309. With kind permission of Elsevier.)

Fig. 5.9 Optical micrographs of samples showing (a) the microstructure near the bottom of wall with fine lamellar Widmanst€atten structure and a small amount of grain boundary α-phase indicated by arrows and (b) the microstructure near the top of wall with variation in lath structures and slightly coarser Widmanst€atten lamellae. The build direction is vertical in both images. (Carroll, B.E., Palmer, T.A., Beese, A.M., 2015. Acta Mater. 87, 309. With kind permission of Elsevier.)

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using the direct energy deposition (DED) method is considered below, in particular, the effect of heat treatment to induce isotropic behavior is tested. For this purpose, the stress-strain curves of the LPD Ti-6Al-4V are compared with Ti-6Al-4V commercial products. Stress-strain curves of commercial Ti-6Al4V are shown in Fig. 5.10. Compare these curves with not heat-treated and heat-treated LPD Ti-6Al-4V curves in Figs. 5.11 and 5.12. Tables 5.4 and 5.5 list the mechanical properties of the commercial LPD Ti-6Al-4V and the heat-treated alloys. The anisotropy observed between the longitudinal and transverse samples in the none heat treated is clearly indicated in Table 5.4. Elongation of the longitudinal samples, for example, is significantly lower than in samples oriented in the transverse direction. The anisotropy is reflected also in the tensile stress and to a lower degree in the yield stress. Heat treatment removes this anisotropy in all mechanical properties—as mentioned earlier—and it can be observed in the listed values in Table 5.5 (L in the tables stands for longitudinal and T means transvers direction).

1200

0.001/s

Eng. Stress(MPa)

1000

800

600

400

China #1 China #2 China #3 USA #1 USA #2 USA #3

200

0 0.00

0.02

0.04

0.06

0.08

0.10

0.12

0.14

0.16

Eng. Strain Fig. 5.10 Commercial products engineering stress-strain curve at 0.001/s. (Woo, S., Lee, Y., Park, L., 2018. EPJ Web Conf. 183, 04002, Open access.)

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1200

2 Layer - 0.001/s

Eng. Stress(MPa)

1000

800

600

400 2L-L #1 2L-L #2 2L-L #3 2L-T #1 2L-T #2 2L-T #3

200

0 0.00

0.02

0.04

0.06

0.08

0.10

0.12

0.14

Eng. Strain

(a) 1200

3 Layer - 0.001/s

Eng. Stress(MPa)

1000

800

600

400 3L-L #1 3L-L #2 3L-L #3 3L-T #1 3L-T #2 3L-T #3

200

0 0.00

0.02

0.04

0.06

0.08

0.10

0.12

Eng. Strain

(b) Fig. 5.11 Engineering stress-strain curve LPD Ti64 alloys according to the direction and layer type at 0.001/s. (a) 2 layer, (b) 3 layer. (Woo, S., Lee, Y., Park, L., 2018. EPJ Web Conf. 183, 04002, Open access.)

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1200

2 Layer - 0.001/s

Eng. Stress(MPa)

1000 800 600 400 2L-LH #1 2L-LH #2 2L-LH #3 2L-TH #1 2L-TH #2 2L-TH #3

200 0 0.00

0.02

0.04

0.06

0.08

0.10

0.12

Eng. Strain

(a) 1200

3 Layer - 0.001/s

Eng. Stress(MPa)

1000 800 600 400 3L-LH #1 3L-LH #2 3L-LH #3 3L-TH #1 3L-TH #2 3L-TH #3

200 0 0.00

0.02

0.04

0.06

0.08

0.10

0.12

Eng. Strain

(b) Fig. 5.12 Engineering stress-strain curve LPD Ti64 alloy with heat treatment at 0.001/s. (a) 2 layer, (b) 3 layer. (Woo, S., Lee, Y., Park, L., 2018. EPJ Web Conf. 183, 04002, Open access.)

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Table 5.4 mechanical properties of commercial and LPD Ti64 alloys at quasi-static regime. Material type

Yield St. (MPa)

Tensile St. (MPa)

Elong. (%)

USA China 2L-L 3L-L 2L-T 3L-T

1043 (11) 1045 (7) 1036 (19) 1033 (34) 1011 (3) 946 (3)

1096 (1) 1071 (9) 1113 (25) 1116 (30) 1086 (7) 1027 (2)

15.3 (0.5) 14.3 (0.3) 7.3 (1.7) 8.3 (0.7) 11.3 (1.3) 11.0 (1.0)

Woo, S., Lee, Y., Park, L., 2018. EPJ Web Conf. 183, 04002.

Table 5.5 mechanical properties of LPD Ti64 alloys with heat-treatment at quasi-static regime. Material type

Yield St. (MPa)

Tensile St. (MPa)

Elong. (%)

2L-LH 3L-LH 2L-TH 3L-TH

889 855 864 877

953 930 934 952

9.7 (1.7) 10.0 (2) 9.7 (1.7) 8.4 (0.5)

(5) (32) (12) (15)

(3) (22) (7) (3)

Woo, S., Lee, Y., Park, L., 2018. EPJ Web Conf. 183, 04002.

5.1.2 In traditionally fabricated Ti-6Al-4V 5.1.2.1 Tensile deformation Tensile test deformation performed at two strain rates and various temperatures up to 1050°C is illustrated in Figs. 5.13 and 5.14. At the lower temperatures (Fig. 5.13), work hardening is observed with very small ductility and the 1500

RT

True stress (MPa)

1250 300°C 1000 750 500°C 500 600°C 250

650°C

0 0

0.5 True plastic strain

1

Fig. 5.13 True stress-true plastic strain curves of a Ti-6Al-4V flat specimens tested between RT and 650°C and a strain rate of 5  104 s1. (Vanderhasten, M., Rabet, L., Verlinden, B., 2008. Mater. Des. 29, 1090. With kind permission of Elsevier.)

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Additive and traditionally manufactured components

True stress (MPa)

150

850°C

750°C 100

800°C 900°C

50

950°C 0 0

1

2

True plastic strain Fig. 5.14 True stress-true plastic strain curves of Ti-6Al-4V flat specimens tested between 750°C and 950°C and a strain rate of 5  104 s1. (Vanderhasten, M., Rabet, L., Verlinden, B., 2008. Mater. Des. 29, 1090. With kind permission of Elsevier.)

fracture strain is limited, but with increasing temperature softening is observed and recovery becomes an important mechanism. As seen in Fig. 5.15, the elongation of the individual grains is parallel to the macroscopic tensile deformation. No recrystallization occurs at these temperatures and no grain boundary sliding was observed. In the 725°C–950°C range and at strain rates of or below 5  103 s1, the main deformation mechanism is grain boundary sliding. Enhanced deformability characterizes the Ti-6Al-4V alloy.

Fig. 5.15 Microstructure of a Ti-6Al-4V sample deformed at 500°C at a strain rate of 5  104 s1 until 30% engineering strain. The tensile direction is horizontal. Grains are stretched in the tensile direction. (Vanderhasten, M., Rabet, L., Verlinden, B., 2008. Mater. Des. 29, 1090. With kind permission of Elsevier.)

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Deformation in AM and traditional manufacturing

At temperatures above 950°C, dynamic grain growth occurs and most likely dislocation creep is the responsible deformation mode. It could be noted that at high-strain rate at temperatures above 725°C, the σ-ε curves are of the steady-state type and hot deformation combined with dynamic recrystallization seems to take place. 5.1.2.2 Compressive deformation In Fig. 5.16, stress-strain curves at two different sizes obtained by compression are illustrated. Strain hardening is seen in the plastic region as commonly observed in metals (see Fig. 5.17). The cylindrical EBM-built specimens

Fig. 5.16 Static true stress-strain curves for different specimen sizes. (Mohammadhosseini, A., Masood, S.H., Fraser, D., Jahedi, M., 2015. Adv. Manuf. 3, 232. With kind permission of Springer Nature.)

1600

Elastic deformation stage Damage fracture stage

1400

Stress (MPa)

1200

B

1000

A

800 600

Plastic deformation stage

400 200 0 0.0

0.1

0.2

0.3

0.4

0.5

Strain (-) Fig. 5.17 Stress-strain curve of Ti-6Al-4V alloy in the quasi-static test. (Hou, X., Liu, Z., Wang, B., Lv, W., Liang, X., Hua, Y., 2018. Materials 11, 938, Open access.)

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Additive and traditionally manufactured components

were tested at low-strain rate of 103 s1. The effect of temperature at various strain rates is illustrated in Fig. 5.18. Stress-strain curves at a constant strain rate of 10,000 s1 in the temperature range 25°C–600°C are illustrated in Fig. 5.19. As would be expected, the level of the stress-strain curves decreases with temperature increase. Also, the strain-hardening rate 2,000

b 2,000

1,500

1,500 Stress (MPa)

Stress (MPa)

a

1,000 -1

4000 s

6000 s-1

500

1,000 4000 s-1 6000 s-1

500

-1

10000 s

10000 s-1

-1

12000 s

0 0.00

0.05

0.10

0.15

0.20

0.25

12000 s-1

0.30

0 0.00

0.35

0.05

0.10

Strain (-)

1,500 Stress (MPa)

1,500 Stress (MPa)

d 2,000

1,000 -1

4000 s

0.10

0.15

0.35

0.20

6000 s-1

10000 s-1

10000 s-1

12000 s-1

12000 s-1

0.25

0.30

0 0.00

0.35

e 2,000

0.05

0.10

0.15 0.20 Strain (-)

0.25

0.30

0.35

f 2,000 4000 s-1

4000 s-1 6000 s-1

1,500

6000 s-1

1,500

10000 s-1 12000 s-1

Stress (MPa)

Stress (MPa)

0.30

4000 s-1

500

Strain (-)

1,000

10000 s-1 12000 s-1

1,000

500

500

0 0.00

0.25

1,000

6000 s-1

500

0.05

0.20

Strain (-)

c 2,000

0 0.00

0.15

0.05

0.10

0.15

0.20

Strain (-)

0.25

0.30

0.35

0 0.00

0.05

0.10

0.15

0.20

0.25

0.30

0.35

Strain (-)

Fig. 5.18 Stress-strain curves of Ti-6Al-4V alloy over a wide range of temperatures of (a) 25°C, (b) 100°C, (c) 200°C, (d) 300°C, (e) 400°C, and (f) 500°C in SHPB test. SHPB stands for split Hopkinson pressure bar. (Hou, X., Liu, Z., Wang, B., Lv, W., Liang, X., Hua, Y., 2018. Materials 11, 938, Open access.)

241

Deformation in AM and traditional manufacturing

2,000

Stress (MPa)

1,500

1,000

500

25 °C 300 °C 600 °C

0 0.00

0.05

0.10

100 °C 400 °C

0.15 0.20 Strain (-)

200 °C 500 °C

0.25

0.30

0.35

Fig. 5.19 Stress-strain curves of Ti-6Al-4V alloy over the wide range of temperature from 25°C to 600°C in SHPB test. (Hou, X., Liu, Z., Wang, B., Lv, W., Liang, X., Hua, Y., 2018. Materials 11, 938, Open access.)

decreases with the increase of temperature as seen in Fig. 5.20, where the ratio of the experimental temperature to room temperature is shown on the abscissa rather than the temperature. In summary, stress-strain curves were illustrated over the temperature range of 25°C–600°C and strain rates 0.0016–12,000 s1 indicating that 1,200

Strain hardening rate (MPa)

1,000 y = 1019.2e-6.58x

800

R2 = 0.9907 600 400 200 0 0

5 10 15 20 25 Ratio of experimental temperature to RT (-)

30

Fig. 5.20 Strain-hardening rate over the wide range of temperature. (Hou, X., Liu, Z., Wang, B., Lv, W., Liang, X., Hua, Y., 2018. Materials 11, 938, Open access.)

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Additive and traditionally manufactured components

strain hardening, strain-hardening rate, and the stress level are inverse functions of the temperature. To confirm the temperature and strain-rate sensitivity of Ti-6Al-4V, additional experiments carried out at high-strain rates and high temperature up to 1100°C can be illustrated (Fig. 5.21). Stress-strain relations can be expressed by the known Ludwik equation given as σ ¼ A + Bεn

(5.2)

where n is the work-hardening coefficient, A is a constant often the yield strength, and B is a material constant. However, good description of the stress-strain relation can be obtained by the Johnson-Cook model given by σ ¼ ðA + Bεn Þð1 + clnε_ ∗ Þ ð1  T ∗m Þ

(5.3)

_ ¼ ε= _ ε_ 0 is the ratio of the test which is illustrated in Fig. 5.22. In Eq. (5.3), ε* strain rate to a reference strain rate, ε is the equivalent plastic strain, T * is the homologous temperature (T  Troom)/(Tmelt  Troom), and A, B, n, c, and m are the material constants. B is the strain-hardening coefficient and n is clearly the strain-hardening exponent. The values are listed in Table 5.6. Fracture after deformation occurred catastrophically at 45° to the compression axis and was associated with adiabatic shear bands under higher strain rates than for the tested cases at all temperatures. Voids are nucleated in 1600 25 °C 1400

True Stress (MPa)

300 °C 1200

500 °C

1000

700 °C

800

900 °C 1100 °C

600 -1

800 s 1400 s-1 1700 s-1 2000 s-1 2500 s-1 3300 s-1

400 200 0

0

0.1

0.2 True Strain

0.3

0.4

Fig. 5.21 Typical true stress-strain curves of Ti-6Al-4V alloy deformed at different strain rates and temperature conditions. (Lee, W.-S., Lin, C.-F., 1998. Mater. Sci. Eng. A 241, 48. With kind permission of Elsevier.)

243

Deformation in AM and traditional manufacturing

1800 Strain rate = 2000 s–1

25 °C

True stress (MPa)

1500

300 °C 1200

500 °C 700 °C

900

900 °C 600

1100 °C measured values calculated values

300 0

0

0.1

0.2 True strain

0.3

0.4

Fig. 5.22 Comparison between predicted and measured stress-strain curves for a strain rate of 2  103 s1. (Lee, W.-S., Lin, C.-F., 1998. Mater. Sci. Eng. A 241, 48. With kind permission of Elsevier.) Table 5.6 Material parameters of Ti–6Al–4V alloy. Materials

A (MPa)

B (MPa)

Τi-6A1-4V

724.7

683.1

n

c

m

0.47

0.035

1.0

Lee, W.-S., Lin, C.-F., 1998. Mater. Sci. Eng. A 241, 48. With kind permission of Elsevier.

the cross section of the shear bands, which then join together to form a crack as seen in Fig. 5.23. The microvoids are nucleated either at weak points in the adiabatic shear zone, or at sections of the shear surfaces, where significant thermal softening has occurred. In summary, the experiments show that flow stress, the material constants, and the work-hardening coefficient are sensitive to strain rate and temperature. The temperature sensitivity increases with the increase of temperature, but is independent of strain rate.

5.2 Deformation in AM Al AA6061 As mentioned in one of the sections above, AA 6061 is a precipitation-hardened alloy based on magnesium and silicon. The interest in this alloy is related to its good mechanical strength. It is extrudable and exhibits good weldability. Because of the mentioned properties, it is very

244

Additive and traditionally manufactured components

Fig. 5.23 Voids and cracks in an adiabatic shear band for the specimen deformed at 700°C under 2  103 s1. (Lee, W.-S., Lin, C.-F., 1998. Mater. Sci. Eng. A 241, 48. With kind permission of Elsevier.)

popular in use for general purposes. It also serves as structural material for construction, transportation, and sports. Various heat treatments result in variation of the microstructure and the consequent change in mechanical properties. The specific treatments are indicated by a letter following the type of alloy; thus, for example, we indicate the annealed alloy as 6061-O (annealed pre-tempered), 6061-T4 (tempered), or 6061-T6 (solutionized and artificially aged). In the following, deformation is discussed in AM and conventional AA6061.

5.2.1 Tensile deformation in Al AA6061 Engineering stress-strain curves were constructed from builds fabricated using Al-AA 6061 tapes in the H-18 condition. (H-18 symbolizes an alloy work hardened to full hardness and not annealed after rolling.) The alloys were obtained using ultrasonic additive manufacturing (UAM), which is a high-strain-rate process. The stress-strain curves tested in the X-, Y-, and Z-directions are compared to the wrought alloy in Fig. 5.24. The specimens loaded along the Z-axis showed a drastic decrease in the strength level and failed in brittle manner. Average strength and ductility values are given in Table 5.7. The wrought alloys were in the T-6 temper condition. As seen in the figure, the UAM specimens loaded in the X- and Y-directions had lower strength and ductility than the wrought alloy with the same composition. A drastic decrease in strength in the Z-direction was observed and

245

Deformation in AM and traditional manufacturing

300

X

Z

Stress(MPa)

250 200 150 100 X Direction Z Direction Bulk properties

50 0 0.00

0.05

0.10 Strain

1 mm

Regions where post deformation characterization was performed

0.15

Fig. 5.24 Typical engineering tensile curves for the UAM-produced specimens in comparison with bulk commercial alloy. Note the drastic decrease of strength along the Z-direction. Areas where post-deformation characterization was performed are also marked in the figure. (Sridharan, N., Gussev, M., Seibert, R., Parish, C., Norfolk, M., Terrani, K., et al., 2016. Acta Mater. 117, 228. With kind permission of Elsevier.) Table 5.7 Mechanical properties of the UAM-specimens and wrought commercial alloy. Type

Wrought alloy X Y Za

Yield stress (MPa)

Ultimate stress (MPa)

Uniform elongation (%)

Total elongation (%)

294

315

7.2

15.4

217 221 46

225 224 –

1.1 0.5 –

5.1 6.0 –

a Fracture stress is given for the Z-direction specimens. Sridharan, N., Gussev, M., Seibert, R., Parish, C., Norfolk, M., Terrani, K., et al., 2016. Acta Mater. 117, 228. With kind permission of Elsevier.

alloys failed in a brittle manner as indicated. The total elongation of the UAM alloys in the X- and T-directions is considerably lower than in the wrought samples (see Table 5.7), but also their strength levels are also much reduced compared to the bulk specimens (wrought). The fractured specimens away from the interface displayed equiaxed dimples, a sign of ductile fracture. This is illustrated in Fig. 5.25c and d. Near the interface, no dimples were seen, which is likely to be the result of delamination at the interface immediately occurring on loading (recall that AM is a layer-by-layer process). No localized plastic deformation can occur in these locations of the build.

246

Additive and traditionally manufactured components

Fig. 5.25 The fractographs of the samples: (a) tested along the “Z”-direction and (b) the ductile regions in (a) at a higher magnification. The presence of dimples at the interface indicative of a ductile failure (c) samples tested along the “X”-direction and (d) samples tested along the “X”-direction [(c) shown at a higher magnification] showing debonding of the tapes at the interface. (Sridharan, N., Gussev, M., Seibert, R., Parish, C., Norfolk, M., Terrani, K., et al., 2016. Acta Mater. 117, 228. With kind permission of Elsevier.)

One can conclude that the mechanical properties in 6061 builds depend on the direction of tests. Loading parallel to the interface, i.e., in the Xdirection lower values of strength and ductility are observed, however, when loading is perpendicular to the interface, namely in the Z-direction brittle failure occurs. The brittle failure is a consequence of strain localization during the UAM process, which is a high-strain-rate deformation as mentioned. High-strain-rate deformation results in strain localization, leading to microvoids at the interface. These microvoids coalesce nucleating cracks that propagate through the interface without macroscopic ductility resulting in fracture in the present case in the Z-direction.

5.2.2 Compressive deformation Stress-strain relationship for the AA6061 specimens obtained by compression is illustrated in Fig. 5.26. The C90-oriented sample shows a higher stress-strain curve. Images of the specimen after compression

247

Deformation in AM and traditional manufacturing

250

Stress [MPa]

200

150 C0 C45

100

C90 50

0 0

10

20

30

40

50

60

70

80

90

Strain [%]

Fig. 5.26 Thesis: characterization of aluminum components produced by additive manufacturing. (Rønneberg, T., Norwegian University of Science and Technology, NTNU Department of Materials Science and Engineering, p. 77. With kind permission of Dr. T. Ronneberg.)

testing are shown in Fig. 5.27. The specimens illustrated in Fig. 5.27 (after compression) show a slight barreling. The flow stress of AA6061 is compared with that of AlSi10Mg in Table 5.8. The stopping criterions are also included in the table and for AA6061 the criterion is the height. Note that the C90-oriented sample has the highest flow stress (as seen also in Fig. 5.26).

5.2.3 Conventional tensile deformation Alloy AA6061 is commonly available in pre-tempered grades such as 6061O (which represents annealed conditions), tempered grades such as 6061-T6 (solutionized and artificially aged) and 6061-T651 (solutionized, stressrelieved stretched, and artificially aged). In Fig. 5.28, elevated temperature 6061-T651 engineering stress-strain relations are shown. As can be seen in Fig. 5.28b, the primary strength reduction occurred from about 250°C down to higher test temperatures leading to decreased yield strength. The strength reduction magnitude is dependent on the isothermal exposure temperature and duration. The yield strength variation with temperature is included among other alloys in Fig. 5.29 and the Young’s modulus variation with temperature of 6061-T651 is seen in Fig. 5.30. The ultimate tensile

248

Additive and traditionally manufactured components

Fig. 5.27 Images of the specimen after compression testing. (a) C0, (b) C45, (c) C90. (Rønneberg, T., Norwegian University of Science and Technology, NTNU Department of Materials Science and Engineering, p. 77. With kind permission of Dr. T. Ronneberg.) Table 5.8 Results of compression tests. AlSi10Mg

AA6061

Specimen

σ fl (MPa)

Stopping criterion

σ fl (MPa)

Stopping criterion

C0 C45 C90

481 466 471

Time Fracture Load

220 226 241

Height Height Height

Rønneberg, T., Norwegian University of Science and Technology, NTNU Department of Materials Science and Engineering, p. 77. With kind permission of Dr. T. Ronneberg.

strength is illustrated also for 6061-T651 in Fig. 5.31. The ductility in terms of the reduction area as a function of temperature is shown in Fig. 5.32. As indicated earlier, the equal channel pressing (ECAP) is an innovative technique, which among its other advantages is a method for grain refinement and thus enables to overcome limitations in material fabrication. In Fig. 5.33,

249

Deformation in AM and traditional manufacturing

a

350 25°C 50°C 100°C 150°C 200°C 250°C 300°C 350°C 400°C 450°C 500°C

Engineering Stress (MPa)

300 250 200 150 100 50 0

b

350

Engineering Stress (MPa)

300 250 200 150

25°C 50°C 100°C 150°C 200°C 250°C 300°C 350°C 400°C 450°C 500°C

100 50 0 0.000

0.001

0.002

0.003

0.004

0.005

Engineering Strain (mm/mm)

Fig. 5.28 6061-T651 engineering stress-strain relations. (Summers, P.T., Chen, Y., Rippe, C.M., Allen, B., Mouritz, A.P., Case, S.W., et al., 2015. Fire Sci. Rev. 4, 1, Open access.)

the change of hardness with temperature of 6061 produced by the ECAP technique is shown. As known alloy 6061 is a precipitation-hardening alloy and as such overaging can occur, which is associated with reduced mechanical properties and also with grain growth. In addition, the number of passes in the ECAP process is an important factor affecting the mechanical properties. Note that the ECAE stands for extrusion performed by ECAP, specifically equal channel angular extrusion. In Fig. 5.34, the variation of the Brinnel hardness

250

Additive and traditionally manufactured components

350 5083-H116 6061-T651 5083-H321 [38] 6061-T651 [38]

Yield Strength (MPa)

300 250 200 150 100 50 0 0

100

200

300

400

500

Temperature (°C)

Fig. 5.29 5083-H116 and 6061-T651 elevated temperature yield strengths (0.2% offset method). Data reported in Kaufman (2000) Ref. [38] is shown for comparison. (Summers, P.T., Chen, Y., Rippe, C.M., Allen, B., Mouritz, A.P., Case, S.W., et al., 2015. Fire Sci. Rev. 4, 1, Open access.)

70

Young’s Modulus (GPa)

60 50 40 30 5083-H116 6061-T651 Eurocode 9 [6] 5083-H321 [38] 6061-T651 [38]

20 10 0 0

100

200

300

400

500

Temperature (°C)

Fig. 5.30 5083-H116 and 6061-T651 elevated temperature Young’s modulus. Data reported in Eurocode 9 (BSI, 2007) and by Kaufman (2000) in Ref. [38] are shown for comparison. (Summers, P.T., Chen, Y., Rippe, C.M., Allen, B., Mouritz, A.P., Case, S.W., et al., 2015. Fire Sci. Rev. 4, 1, Open access.)

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Deformation in AM and traditional manufacturing

400 5083-H116 6061-T651 5083-H321 [38] 6061-T651 [38]

Ultimate Tensile Strength (MPa)

350 300 250 200 150 100 50 0 0

100

200

300

400

500

Temperature (°C)

Fig. 5.31 5083-H116 and 6061-T651 elevated temperature ultimate strengths. Data reported in Kaufmann et al. (1999) in Ref. [38] are shown for comparison. (Summers, P.T., Chen, Y., Rippe, C.M., Allen, B., Mouritz, A.P., Case, S.W., et al., 2015. Fire Sci. Rev. 4, 1, Open access.) 100 5083-H116 6061-T651 Reduction of Area (%)

80

60

40

20

0 0

100

200

300

400

500

Temperature (°C)

Fig. 5.32 5083-H116 and 6061-T651 reduction in area after failure during tension testing at elevated temperatures (Allen, 2012). (Summers, P.T., Chen, Y., Rippe, C.M., Allen, B., Mouritz, A.P., Case, S.W., et al., 2015. Fire Sci. Rev. 4, 1, Open access.)

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Additive and traditionally manufactured components

BRINELL HARDNESS

160 110C, 4 PASSES 110C, 2 PASSES 110C, 8 PASSES 110C, 4 PASSES 110C, 2 PASSES 110C, 8 PASSES

120 b

peak aged

80 (ECAE) 40 c overaged 0 0

100

200

300

400

ANNEALING TEMPERATURE DEG C

Fig. 5.33 The effect of number of ECAE passes at 110°C and temperature of postdeformation isochronal annealing (1 h) on the hardness of AA 6061. (Ferrasse, S., Segal, V.M., Hartwig, K.T., Goforth, R.E., 1997. J. Mater. Res. 12, 1253. With kind permission of Cambridge University Press.)

BRINELL HARDNESS

160 110C, 4 PASSES 170C, 4 PASSES 140C, 4 PASSES 110C, 4 PASSES 140C, 4 PASSES 170C, 4 PASSES

120 b

peak aged

80

(ECAE) 40 c

overaged

0 0

100

200

300

400

ANNEALING TEMPERATURE DEG C

Fig. 5.34 The effect of deformation temperature after four ECAE passes and temperature of post-deformation annealing on the hardness of AA 6061. (Ferrasse, S., Segal, V.M., Hartwig, K.T., Goforth, R.E., 1997. J. Mater. Res. 12, 1253. With kind permission of Cambridge University Press.)

253

Deformation in AM and traditional manufacturing

with temperature is shown only for four passes, thus the variation in the number of passes is eliminated and only the effect of temperature is shown for the aged and overaged AA6061 alloy (Fig. 5.35). The microstructures obtained by optical microscopy of the AA6061 alloys by the four-pass ECAE process treatment are shown in Fig. 5.36. The peak-aged AA6061 had a grain size of 200– 500 μm as illustrated in Fig. 4.36a and b. After four ECAE passes, the strain state is uniform and is characterized by the presence of very thin parallel and elongated flow lines or zones, where the structure has been sheared to a similar state of deformation. Static recrystallization begins near 300°C and is fully achieved at approximately 400°C with new grains. The grains in the overaged material are somewhat larger than those of the peak-aged material (see Fig. 5.36d and f ). TEM micrographs of the peak-aged AA6061 are shown in Fig. 5.37. Fragmented structure is produced after four ECAE passes as seen in Fig. 5.37a and b, which have large-angle misorientation (revealed by XRD). Elongated, fine grains with an average grain size of 0.2 μm are produced by the ECAE, which produces submicron-grained structure in AA6061. Annealing at 300°C produces grain growth up to
0.8–1 μm. The TEM of overaged AA6061 is shown in Fig. 5.38. Large particles >1 μm of different shape, size, and orientation are seen in Fig. 5.38a representing overaged specimens after deformation. Far away from the large second-phase particles, the structure is quite homogeneous as seen in

UTS (MPa)

300

overaged

80 60 40

200 peak aged

100

20

ELONGATION (%)

peak aged

400

overaged 0 0

100 200 300 ANNEALING TEMPERATURE DEG C

0 400

Fig. 5.35 The effect of isochronal (1 h) annealing temperature on ultimate tensile strength (○,●) and relative elongation (□,■) of the processed peak-aged (○,□) and overaged (●,■) AA 6061 after four passes at 110°C. (Ferrasse, S., Segal, V.M., Hartwig, K.T., Goforth, R.E., 1997. J. Mater. Res. 12, 1253. With kind permission of Cambridge University Press.)

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Additive and traditionally manufactured components

(a)

200 mm

(b)

200 mm

(c)

50 mm

(d)

50 mm

(e)

50 mm

(f)

50 mm

Fig. 5.36 Microstructures of AA6061 observed by optical microscopy; (a) peak-aged material before ECAE deformation; (b) overaged material before ECAE deformation; (c) peak-aged material after deformation (four passes at 110°C) and annealing at 300°C, 1 h; (d) overaged material after deformation (four passes at 110°C) and annealing at 300°C, 1 h; (e) peak-aged material after deformation (four passes at 110°C) and annealing at 400°C, 1 h; and (f) overaged material after deformation (four passes at 110°C) and annealing at 400°C, 1 h. (Ferrasse, S., Segal, V.M., Hartwig, K.T., Goforth, R.E., 1997. J. Mater. Res. 12, 1253. With kind permission of Cambridge University Press.)

Fig. 5.38b. The grain growth during annealing is particularly important at 210°C, the grain size measures around 0.4 μm (Fig. 5.38c) and increases at 300°C to 1–2 μm. Thus, ECAE is an efficient technique for grain refinement in 6061 aluminum alloys. Pre-extrusion heat treatment controls the size of second-phase particles. Due to the deformation at relatively low temperatures

255

Deformation in AM and traditional manufacturing

A

(a)

200 nm

(c)

(b)

200 nm

200 nm

Fig. 5.37 TEM photomicrograph of overaged AA 6061: (a, b) As-deformed overaged material (four passes at 110°C); (c) overaged material after deformation (four passes at 110°C) and annealing at 210°C, 1 h. (Ferrasse, S., Segal, V.M., Hartwig, K.T., Goforth, R.E., 1997. J. Mater. Res. 12, 1253. With kind permission of Cambridge University Press.)

grain size of 0.2–0.4 μm is obtained, while in post-extrusion annealing recrystallization results in 5–15 μm grain size. Strong and ductile, uniform, and stable second-phase small particles are produced at a peak-aged pre-extrusion treatment. As the post-extrusion annealing approaches the temperature of overaging, large particles are gradually formed by the agglomeration of the small particles. Dynamic recrystallization at ECAE is promoted by large particles during pre-extrusion overaging, but also high levels of strain

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Additive and traditionally manufactured components

(a)

200 nm

(c)

(b)

100 nm

100 nm

Fig. 5.38 TEM photomicrograph of overaged AA6061: (a, b) As-deformed overaged material (four passes at 110°C); (c) overaged material after deformation (four passes at 110°C) and annealing at 210°C, 1 h. (Ferrasse, S., Segal, V.M., Hartwig, K.T., Goforth, R.E., 1997. J. Mater. Res. 12, 1253. With kind permission of Cambridge University Press.)

nonuniformity and high internal stresses develop around the large particles. Varying the ECAE parameters and the conditions of pre- and post-extrusion heat treatment provides flexibility for controlling dynamic recrystallization to obtain specific structure and mechanical properties.

5.2.4 Conventional compressive deformation Split Hopkinson pressure bar (SHPB) is used to characterize materials subjected to dynamic loads at different strain rates. SHPB can be applied to

257

Deformation in AM and traditional manufacturing

compressive or tensile uniaxial load. The method was also used to obtain true stress-strain curves by compression. True stress vs true strain and true stress vs true strain rate are illustrated in Figs. 5.39 and 5.40. It would be of interest to show the compression data on a new alloy based on

Fig. 5.39 Compression tests results: true stress vs true strain. (Mancini, E., Sasso, M., Rossi, M., Chiappini, G., Newaz, G., Amodio, D., 2015. J. Dyn. Behav. Mater. 1, 201. With kind permission of Springer Nature.)

a 5000 4500 4000 True Strain Rate [1/s]

3500 3000 Test 1 - 140 1/s

2500 Test 2 - 300 1/s

2000 Test 3 - 380 1/s

1500 Test 4 - 660 1/s

1000 Test 5 - 1250 1/s

500

Test 6 - 1920 1/s

0 0.0

0.2

0.4

0.6 0.8 True Strain [mm/mm]

1.0

1.2

1.4

Fig. 5.40 Compression tests results: true strain rate vs true strain. (Mancini, E., Sasso, M., Rossi, M., Chiappini, G., Newaz, G., Amodio, D., 2015. J. Dyn. Behav. Mater. 1, 201. With kind permission of Springer Nature.)

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Additive and traditionally manufactured components

AA6061, the difference being in somewhat higher Si, Mg, and Cu contents. For comparison, the percent of these additives in the conventional AA6061 and the new alloy based on AA6061 are indicated below: in AA6061 the percent of the added elements are: Si-0.4%–0.8%, Mg-0.8%–1.2%, and Cu-0.15%–0.4%, while in the new alloy they are: Si-0.9%, Mg-1.2%, and Cu-0.6%. True stress-true strain curves by hot compression of the new alloy are shown in Fig. 5.41 for different temperatures in the temperature range 350°C–550°C. Compression true stress-true strain curves at various strain rates (temperature range 350°C–550°C and strain-rate range 0.005–5 s1) were carried out. Optical microstructures of the deformed specimens under different conditions are shown in Fig. 5.42. Elongated grains can be seen obtained at different strain rates at the temperatures of the deformation. TEM images

(b)

(a) 70

70

60

350°C

50

True stress, MPa

True stress, MPa

80

350°C

40 30 20

400°C 450°C 500°C 550°C

10 0 0.0

0.1

0.2

0.3

0.4

0.5

0.6

0.7

0.8

60 50 40 400°C 30 450°C

20

500°C

10

550°C 0 0.0

0.9

0.1

0.2

0.3

True strain 80

(d) 100

70

90

0.6

0.7

80

350°C

60 50

400°C

40 450°C

30

500°C

20

550°C

0.8

0.9

350°C

70 60

400°C

50 450°C

40

500°C

30

550°C

20

10 0

0.5

True strain

True stress, MPa

True stress, MPa

(c)

0.4

10 0.0

0.1

0.2

0.3

0.4

0.5

True strain

0.6

0.7

0.8

0.9

0 0.0

0.1

0.2

0.3

0.4

0.5

0.6

0.7

0.8

0.9

True strain

Fig. 5.41 True stress-true strain curves of the new Al-Mg-Si-Cu alloy during hot compression deformation. (a) ε_ ¼ 0.005 s1, (b) ε_ ¼ 0.05 s1, (c) ε_ ¼ 0.5 s1, and (d) ε_ ¼5 s1. (Zhang, H., Li, L., Yuan, D., Peng, D., 2007. Mater. Charact. 58, 168. With kind permission of Elsevier.)

259

Deformation in AM and traditional manufacturing

(a)

(b)

50mm

(c)

50mm

(d)

100mm

100mm

Fig. 5.42 Optical deformed microstructures of the new Al-Mg-Si-Cu alloy under different conditions. (a) T ¼ 350°C, ε_ ¼ 0.05 s1; (b) T ¼ 350°C, ε_ ¼ 5 s1; (c) T ¼ 550°C, ε_ ¼ 0.05 s1; and (d) T ¼ 550°C, ε_ ¼ 5 s1. (Zhang, H., Li, L., Yuan, D., Peng, D., 2007. Mater. Charact. 58, 168. With kind permission of Elsevier.)

also at different strain rates and at various temperatures are seen in Fig. 5.43. Actually, the graphs seen in Fig. 5.41 are hot deformation curves. The following equations are commonly applied:
 Q ε_ ¼ A σ exp  RT
 Q 00 ε_ ¼ A exp ðβσ Þ exp  RT
 Q _ε ¼ Að sinh ασ Þn exp  RT
 Q Z ¼ ε_ exp  RT 0 n0

(5.4) (5.5) (5.6) (5.7)

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Additive and traditionally manufactured components

Fig. 5.43 Transmission electron micrographs: (a) 550°C, 5 s1, substructure with higher density dislocations 120° grain boundary; (b) 450°C, 5 s1, subgrain size becomes finer; (c) 450°C, 0.05 s1, little dynamic precipitation and some coalesced particles exhibited on the grain boundaries; and (d) 550°C, 5 s1, very finer precipitates in the subgrain. (a) T ¼ 550°C, ε_ ¼5 s1; (b) T ¼ 450°C, ε_ ¼5 s1; (c) T ¼ 450°C, ε_ ¼ 0.05 s1; and (d) T ¼ 550°C, ε_ ¼5 s1. (Zhang, H., Li, L., Yuan, D., Peng, D., 2007. Mater. Charact. 58, 168. With kind permission of Elsevier.)

A, A0 , A00 , n, n0 , α, and β (¼ αn) are constants and Q is the activation energy for deformation and Z is the Zener-Hollomon parameter. It was found that the power law (Eq. 5.1), and the exponential law (Eq. 5.2) break at high stress and at a low stress, respectively. The hyperbolic Eq. (5.3) is more general and suitable for stresses over a wide range. The peak stress is plotted vs stress for various strain rates in Fig. 5.44a and vs the inverse temperature in

261

Deformation in AM and traditional manufacturing

(a)

(b) 3 2 1

Ine×

0

1.T=350°C 2.T=400°C 3.T=450°C 4.T=500°C 5.T=550°C

5

4

3

1.2

4

0.8

3 2 1

1

2

In [sinh(as)]

4

-1 -2 -3

0.4 0.0 -0.4

1.e× = 0.005s-1 2.e× = 0.005s-1

-0.8

3.e× = 0.5s-1 4.e× = 5s-1

-4 -5 -6 -7 10

20

30

40

50

s

60

70

80

90

100

-1.2

0.0012

0.0013

0.0014 0.0015 1/T, K-1

0.0016

0.0017

Fig. 5.44 Relationships between peak stress with: (a) strain rate and (b) deformation temperature. (Zhang, H., Li, L., Yuan, D., Peng, D., 2007. Mater. Charact. 58, 168. With kind permission of Elsevier.)

Fig. 5.44b, respectively. The activation energy for hot deformation, derived from the Arrhenius relation of Fig. 5.44b is 236 kJ/mol. In conclusion, we can summarize, the true stress-true strain behavior at small strains σ3 σ3 > σ2

σ2 > σ1

γ

σ1

Int

Fig. 7.7 Logarithmic creep. The lines are shown for different stresses. (Pelleg, J., 2017. Creep in Ceramics. Springer.)

Fig. 7.8 Larson-Miller parameters (PLM) as a function of creep strain, stress, and alloy. (Bush, R.W., Brice, C.A., 2012. Mater. Sci. Eng. A 554, 12. With kind permission of Elsevier.)

Time-dependent deformation creep

385

As expected, the plot indicates that the Larson-Miller parameter increases with increasing creep strain and decreasing stress. Also, the erbiumcontaining Ti alloy is included. Recall that the Larson-Miller parameter (LMP) is a method to predict the lifetime of a material versus time and temperature based on the Arrhenius rate equation and is expressed as P ðLM Þ ¼ T ð logt + C Þ

(7.16)

7.3 Compressive creep in AM Ti6Al4V AM of Ti6Al4V alloy by selective laser melting (SLM) technique is presented here for the evaluation of its creep resistance by compressive test. Specimens were heat treated and compared with as-fabricated Ti6Al4V. Optical micrographs before and after heat treatment of Ti6Al4V is shown in Fig. 7.9. As can be seen the grains are columnar along the building direction. It is the consequence of the heat discharge in a single direction during solidification. This is generally observed in metals during AM. SEM (scanning electron microscope) images of etched specimens before and after heat

Fig. 7.9 Optical micrographs illustrating the columnar grains of (a) as-fabricated specimen and (b) heat-treated specimen. (Kim, Y.-K., Park, S.-H., Yu, J.-H., AlMangour, B., Lee, K.-A., 2018. Mater. Sci. Eng. A 715, 33. With kind permission of Elsevier.)

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Additive and traditionally manufactured components

treatment are shown in Fig. 7.10. In the as-fabricated specimen (Fig. 7.10a), a fully martensitic structure is seen while in the heat-treated specimen Widmanstatten structure is observed (Fig. 7.10b) with α and β coexisting. The β phase is located in the boundaries of the α colonies. Equilibrium phases at room temperature consisted of about 95% α phase and 5% β phase. Electron back-scattering diffraction (EBSD) was done on polished specimens for further analysis and is illustrated in Fig. 7.11. Compression creep curves for the as-fabricated and the heat-treated specimens are illustrated in Fig. 7.12. No tertiary creep is shown and the creep curve observed had two stages: transient creep and steady-state creep (see Fig. 7.1a). Recall that the steady-state creep is a balance between hardening and dynamic deformation.

Fig. 7.10 SEM images in the SE mode after chemical etching of the microstructures of (a) as-fabricated specimen and (b) heat-treated specimen. (Kim, Y.-K., Park, S.-H., Yu, J.-H., AlMangour, B., Lee, K.-A., 2018. Mater. Sci. Eng. A 715, 33. With kind permission of Elsevier.)

387

Time-dependent deformation creep

(a)

(b)

80 um

(c)

3 mm

80 um

Fig. 7.11 EBSD inverse pole figure maps of SLM-created Ti-6Al-4V alloy: (a) as-fabricated specimen, (b) heat-treated specimen, and (c) high-magnification phase map of heattreated specimen (red—β phase, green—α phase). (Kim, Y.-K., Park, S.-H., Yu, J.-H., AlMangour, B., Lee, K.-A., 2018. Mater. Sci. Eng. A 715, 33. With kind permission of Elsevier.)

0.08

(b) 0.10

Stress [MPa] 140 120 100 80

Creep strain

Creep strain

(a) 0.10

0.06 0.04 0.02

0.08

Stress [MPa] 140 120 100 80

0.06 0.04 0.02

0.00

0.00 0

40000 80000 120000 160000

0

40000 80000 120000 160000

Time [s]

Time [s]

Fig. 7.12 Compressive creep curves of (a) as-fabricated specimen and (b) heat-treated specimen at 500°C. (Kim, Y.-K., Park, S.-H., Yu, J.-H., AlMangour, B., Lee, K.-A., 2018. Mater. Sci. Eng. A 715, 33. With kind permission of Elsevier.)

The heat-treated curves are at a lower level on the creep strain-time curve. As a result, the primary creep (transient creep) extent of the as-fabricated samples is larger than the heat-treated ones at the same stress. Amin et al. (1970) suggested a relation for creep in terms of strain as ε ¼ ε0 + ε_ s t + εt ð1  exp ðrt ÞÞ

(7.17)

here εt is the limiting transient strain (namely the total transient strain) and r is the ratio of the transient creep rate to the transient creep strain, ε is the amount of creep strain, ε0 refers to the instantaneous strain and ε_ is the

388

Additive and traditionally manufactured components

steady-state creep rate. The steady-state creep rate according to SherbyDorn was presented in Eq. (7.10), and rewritten here as   Q n ε_ ¼ Aσ exp (7.18) RT where n is the stress exponent. A low steady-state creep in the heat-treated Ti6Al4V compared to the as-fabricated is shown in Fig. 7.13. The microstructure after creep deformation for the as-fabricated and heat-treated specimens is illustrated in Fig. 7.14.

7.4 Tensile creep in CP Ti6Al4V As already indicated earlier any material, polymer, metal, or ceramic can undergo time-dependent deformation under constant stress if exposed to some temperature for a certain length of time, which depends on the material. The rate of deformation is a function of exposure time, temperature, and applied load and depends on material properties. Creep is the tendency of a solid material to move slowly under the influence of mechanical stress but still below the yield strength and deform permanently. Creep is more severe in materials that are subjected to heat for long periods and

Steady state creep rate [s-1]

10-6

As-fabricated Heat-treated

n = 2.05

10-7

n = 3.12

10-8

80

90

100

110

120

130

140

Applied stress [MPa]

Fig. 7.13 Stress dependences of the steady-state creep rate of the as-fabricated and heat-treated specimens. (Kim, Y.-K., Park, S.-H., Yu, J.-H., AlMangour, B., Lee, K.-A., 2018. Mater. Sci. Eng. A 715, 33. With kind permission of Elsevier.)

Time-dependent deformation creep

389

Fig. 7.14 Microstructures after creep deformation of SLM-created Ti64 alloy: (a) asfabricated specimen and (b) heat-treated specimen. (Kim, Y.-K., Park, S.-H., Yu, J.-H., AlMangour, B., Lee, K.-A., 2018. Mater. Sci. Eng. A 715, 33. With kind permission of Elsevier.)

generally increase as their temperature increase until a critical failure (dimensional and shape changes, which makes a component useless) is reached near the melting point. Therefore, creep is of concern to engineers when designing components to operate under high stress or high temperatures. However, creep may occur at relatively moderate temperatures in many cases. Even some ceramics with low-temperature ductility may creep 0.5 Tm. An important consideration is the choice of the type of material, namely, single or polycrystalline material. In order to eliminate grain boundary sliding in the material while exposed to some elevated temperature single crystal is the obvious choice.

390

Additive and traditionally manufactured components

However, for many reasons such as other mechanical properties or cost factor consideration, polycrystalline materials are used despite their limitations for high-temperature applications. Grain size is a factor. A compromise must be made between small grain size, which enhances most of the mechanical properties, and large grained material, which is preferred for long-time exposure (creep) at high temperature. Ti6Al4V alloy also known as Ti64 is the most commonly used titanium alloy. It exhibits good mechanical and physical properties and due to its lightweight it is used in aerospace applications. These unique properties are strongly affected by chemical composition, microstructure, deformation, and heat-treatment history. The most commonly used titanium alloy is the two-phase (α + β) alloy. Significant research was devoted to improving the microstructure and attaining appropriate mechanical properties. It has furthermore been shown that the cooling rate seems to be one of the most important parameters affecting microstructural development whereas slow and intermediate cooling rates lead to the nucleation and growth of α-lamellae into initial β grains through a diffusion-controlled process. The higher cooling rates induce a martensitic transformation. The optical micrograph of conventionally produced (CP) Ti6Al4V is shown in Fig. 7.15. The Ti6Al4V alloy is a two-phase Widmanst€atten structure (average grain size 395 μm) where the α colonies are lamellar (parallel oriented lamellae) with lath size (width) of about 3.2–4.0 μm. Creep curves of Ti6Al4V at 500°C and 600°C (but under different stresses) are illustrated in Figs. 7.16 and 7.17. The curves of the lamellar structure shown in Fig. 7.15 exhibit well-defined primary (transient), secondary (steady state), and tertiary (accelerated creep) stages. A reduction in primary creep extent (period) with increase in stress and temperature is shown in Figs. 7.16 and 7.17. The increase in tertiary creep rate is associated with the development of necking and nucleation of microvoids and their coalescence. Further at this stage there is also an increase in the mobile dislocation density. The major part of the creep (see Figs. 7.16 and 7.17) dominating the life span of Ti6Al4V is the constant creep rate where hardening and recovery processes compete with each other, thus determining the steady-state creep. The variation in the steady-state creep rate with applied stress in Ti6Al4V is shown in Fig. 7.18. As shown in this figure the steady-state creep level is higher at the higher temperature. Probably dislocation climb is activated as a result of the exposure to the high temperature. The Sherby-Dorn expression for steady-state creep indicated in Eq. (7.18) shows its combined dependence on stress and temperature. The stress exponents of n ¼ 11.3 and n ¼ 5.2 are obtained from the slope of the curves described by Eq. (7.18). The activation energy was calculated to be 415 kJ/mol at 291 MPa.

391

Time-dependent deformation creep

Fig. 7.15 Optical micrographs of Ti-6Al-4V after heat treatment for 0.5 h at 1050°C followed by furnace cooling to 700°C for 1 h and air cooled to room temperature showing typical features of the (a) grain morphology and (b) lamellar structure. (Barboza, M.J. R., Perez, E.A.C., Medeiros, M.M., Reis, D.A.P., Nono, M.C.A., Neto, F.P., Silva, C.R.M., 2006. Mater. Sci. Eng. A 428, 319. With kind permission of Elsevier.)

These values of n and Q are higher than those of Ti. An explanation is given by considering that creep occurs under a reduced stress of (σ  σ 0) and the stress and temperature dependence can be written as ε_ s

¼ A∗ ð σ  σ

  Qs∗ 0 Þ exp  RT p

(7.19)

where A* is a constant, p  ¼4, and Q∗s is the creep activaton energy for self or atom diffusion. σ 0 is referred to as the threshold stress. Assuming that

392

Additive and traditionally manufactured components

(a)

1.2

Ti-6AI-4V 1.0

T=500°C

Creep Strain [%]

σ=291 MPa 0.8 0.6

0.4

0.2 Creep test interrupted at 5.94 x 105s 0.0 0

100

300

200

400

500

600

Time [103s]

7

(b) Ti-6AI-4V

6

T=500°C σ=403 MPa

Creep Strain [%]

5 4 3 2 1 0

0

10

20

30

40

50

60

70

80

90

100

Time [103s]

Fig. 7.16 Typical creep curves of Ti-6Al-4V at 500°C: (a) 291 MPa and (b) 403 MPa. (Barboza, M.J.R., Perez, E.A.C., Medeiros, M.M., Reis, D.A.P., Nono, M.C.A., Neto, F.P., Silva, C.R.M., 2006. Mater. Sci. Eng. A 428, 319. With kind permission of Elsevier.)

the stress threshold leads to high stress exponent values, the magnitude of its values can be found from a plot on linear scale of the experimental values of ε_ 1=4:3 and σ and extrapolating linearly to ε_ 1=4:3 ¼ 0 as illustrated in Fig. 7.19. The stress hold stresses at 500°C and 600°C are 217.81 and

393

Time-dependent deformation creep

(a)

6

Ti-6AI-4V 5

T=600°C

Creep Strain [%]

σ=97 MPa 4

3

2

1 Creep test interrupted at 2.016 × 106s 0 0

200

400

600

800 1000 1200 1400 1600 1800 2000 2200 Time [103s]

18

(b)

16

Ti-6AI-4V T=600°C

Creep Strain [%]

14

σ=208 MPa

12 10 8 6 4 2 0

0

10

20

30

40

50

60

70

3

Time [10 s]

Fig. 7.17 Typical creep curves of Ti-6Al-4V at 600°C: (a) 97 MPa and (b) 208 MPa. (Barboza, M.J.R., Perez, E.A.C., Medeiros, M.M., Reis, D.A.P., Nono, M.C.A., Neto, F.P., Silva, C.R.M., 2006. Mater. Sci. Eng. A 428, 319. With kind permission of Elsevier.)

34.47 MPa, respectively. As shown in the figure σ 0 decreases with increasing temperature which can be associated with the largest thermal activation energy and dislocation mobility. The true stress exponents for creep thus obtained are p ¼ 4.4 at 500°C and p ¼ 4.1 at 600°C. The creep is

394

Additive and traditionally manufactured components

Fig. 7.18 Dependence of steady-state creep rate on applied stress at 500°C and 600°C. (Barboza, M.J.R., Perez, E.A.C., Medeiros, M.M., Reis, D.A.P., Nono, M.C.A., Neto, F.P., Silva, C. R.M., 2006. Mater. Sci. Eng. A 428, 319. With kind permission of Elsevier.)

Fig. 7.19 The plot of ð_ε Þ1=4:3 against σ for Ti-6Al-4V (the extrapolation of the linear regression to zero creep rate gives the threshold stress). (Barboza, M.J.R., Perez, E.A.C., Medeiros, M.M., Reis, D.A.P., Nono, M.C.A., Neto, F.P., Silva, C.R.M., 2006. Mater. Sci. Eng. A 428, 319. With kind permission of Elsevier.)

Time-dependent deformation creep

395

Fig. 7.20 Plot of steady-state creep rate against effective stress for Ti and Ti-6Al-4V at 500°C and 600°C (present work). (From Zhu, S.J., Mukherji, D., Chen, W., Lu, Y.X., Wang, Z. G., Wahi, R.P., 1998. Mater. Sci. Eng. A 256, 301; Barboza, M.J.R., Perez, E.A.C., Medeiros, M.M., Reis, D.A.P., Nono, M.C.A., Neto, F.P., Silva, C.R.M., 2006. Mater. Sci. Eng. A 428, 319. With kind permission of Elsevier.)

consistent with diffusion-controlled dislocation climb process. For comparison, the steady-state creep versus the effective stress of Ti6Al4V and Ti is plotted in Fig. 7.20. Thus, by introducing threshold stress, stress exponents of 4 can be obtained consistent with lattice diffusion-controlled climb in α-Ti.

7.5 Compressive creep in CP Ti6Al4V A well-known observation is that large grain materials are more creep resistant than fine-grained ones. Therefore, it might be of interest to consider the behavior of fine-grained Ti6Al4V. Uniaxial constant stress compression tests were performed at 648–698 K and at stresses between 300 and 600 MPa on the ultrafine-grained (UFG) processed alloy and, for comparison purposes, on its coarse-grained (CG) state. The material for the creep tests was prepared by multiaxial forging. The severe plastic deformation (SPD) by forging was performed to significantly strengthen the material at ambient temperatures (and induce very high tensile ductilities at elevated temperatures).

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Additive and traditionally manufactured components

Fig. 7.21 Microstructure of UFG Ti alloy processed by multiaxial forging. (Král, P., Dvorak, J., Zherebtsov, S., Salishchev, G., Kvapilová, M., Skleniċka, V., 2013. J. Mater. Sci. 48, 4789. With kind permission of Springer Nature.)

Due to grain growth and recrystallization (at about 0.3 Tm), UFG materials are tested at limited temperatures in the range 0.2–0.5 Tm. Homogeneous structure was obtained after forging (grain size 150 nm) as shown in Fig. 7.21. Inside the grains, dislocation and nonhomogeneous diffraction contrast are observed. Annealing at 773 K for 1 h induced some grain growth resulting in a grain size of 215 nm and the structure is inhomogeneous with some equiaxed and large elongated grains as illustrated in Fig. 7.22. The stress-dependent minimum creep rate is shown in Fig. 7.23. As can be seen fine-grained Ti6Al4V exhibited much faster minimum creep rate than CG ones, which is expected as indicated in the beginning of this section. The stress exponent, n, decreases from 17 for CG material to 4 for the UFG Ti6Al4V. The stress exponent, n, of 4–5 for UFG Ti6Al4V indicates a creep mechanism associated with an intergranular dislocation process involving dislocation glide and climb. The activation energy for the minimum creep rate measured at a stress of 300 MPa and in the temperature range 648–698 K for UFG is shown in Fig. 7.24. The value of activation energy for UFG alloy is 375 kJ/mol. The minimum creep rate measured at 673 K changes significantly with decreasing applied stress. Also, the strain rate to fracture, σ f, varies with decreasing applied stress. As shown in Fig. 7.25, the strain rate to fracture and the

397

Time-dependent deformation creep

1010

0001

2110

1 mm

Fig. 7.22 Microstructure of UFG Ti alloy annealed at 773 K for 1 h. (Král, P., Dvorak, J., Zherebtsov, S., Salishchev, G., Kvapilová, M., Skleniċka, V., 2013. J. Mater. Sci. 48, 4789. With kind permission of Springer Nature.)

. MINIMUM CREEP RATE εmin [S-1]

10-5

10-6

10-7

n=4

coarse-grained state UFG state

n=16.7

10-8

TAV 673 K, 300-600MPa Tension d=2ws

-9

10

100

1000

STRESS s [MPa]

Fig. 7.23 Stress dependences of minimum creep rate measured at 673 K for unpressed CG and UFG Ti-6Al-4V alloys. (Král, P., Dvorak, J., Zherebtsov, S., Salishchev, G., Kvapilová, M., Skleniċka, V., 2013. J. Mater. Sci. 48, 4789. With kind permission of Springer Nature.)

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Additive and traditionally manufactured components

Fig. 7.24 Activation energy of creep Qc determined for UFG alloy. (Král, P., Dvorak, J., Zherebtsov, S., Salishchev, G., Kvapilová, M., Skleniċka, V., 2013. J. Mater. Sci. 48, 4789. With kind permission of Springer Nature.)

. MINIMUM CREEP RATE emin [s–1] -7

10-6

10

10-5

STRAIN TO FRACTURE [%]

80 UFG Ti-6Al-4V alloy creep 673K 300-600MPa

60

40

20 stress min. creep rate 0 200

300

400

500

600

STRESS [MPa]

Fig. 7.25 Changes in strain to fracture with applied stress and minimum creep rate measured at 673 K. (Král, P., Dvorak, J., Zherebtsov, S., Salishchev, G., Kvapilová, M., Skleniċka, V., 2013. J. Mater. Sci. 48, 4789. With kind permission of Springer Nature.)

399

Time-dependent deformation creep

minimum creep rate decrease with increasing stress. Following creep exposure at 673 K under an applied stress of 300 MPa, the grains of about 180 nm and containing about 84% high angle grain boundaries were randomly oriented. The grain structure was bimodal with large elongated grains randomly distributed and less equiaxed grains. The larger grains in the bimodal structure can improve the deformation behavior of UFG Ti6Al4V by plastic deformation inside the grains. The microstructure after exposure to creep is shown in Fig. 7.26. In summary, creep at 673 K in UFG samples exhibited faster minimal creep rate by 1–2 orders of magnitude than in CG Ti6Al4V. The difference in the minimal creep rates between the UFG and the CG is significantly stress dependent. The observations are consistent with the creep controlled by the recovery of dislocations at high-angle grain boundaries where the dislocations are stored. The activation energy for creep of the UFG structure samples tested are much higher than that for the selfdiffusion of Ti or the solute diffusion of Al in α-Ti. It is likely that one of the key factors in the creep damage is the degradation near grain boundaries because cavitation at grain boundaries and crack nucleation lead to intergranular fracture.

1010

0001

2110

1 mm

Fig. 7.26 Microstructure of UFG Ti alloy after subsequent creep exposure at 673 K under 300 MPa. (Král, P., Dvorak, J., Zherebtsov, S., Salishchev, G., Kvapilová, M., Skleniċka, V., 2013. J. Mater. Sci. 48, 4789. With kind permission of Springer Nature.)

400

Additive and traditionally manufactured components

7.6 Tensile creep in AM Al10SiMg AM such as direct metal laser sintering (DMLS) has been primarily limited to Al10SiMg the reason is likely to be its low coefficient of thermal expansion (CTE) due to the presence of Si, which enhances processability. Al 6061 is a commonly used alloy in the industry; however, its higher CTE presents greater difficulties in processing. Controlling shrinkage-induced warping is problematic and AA6061 is not widely used in the AM industry (Fulcher et al., n.d.). The use of AA6061 still remains a challenge and needs to be developed for AM applications. Therefore, this section considers Al10SiMg instead of AA6061. Conventional presentation of creep curves as creep strain against time for Al10SiMg is illustrated in Fig. 7.27 at various temperatures under a 117 MPa stress. Creep curves at 225°C under various stresses are shown in Fig. 7.28. Figs. 7.29 and 7.30 show the variation of ln (strain rate) versus the inverse temperature and ln (strain rate) versus ln (stress), respectively. The slope of Fig. 7.29 is used to determine the activation energy, while from the slope of Fig. 7.30 the stress exponent can be derived. These plots are linear and the equation used for their analysis is Eq. (7.18).

Fig. 7.27 Creep curves of SLM AlSi10Mg specimens at various temperatures under 117 MPa stress. SLM stands for selective laser melting. (Uzan, N.E., Shneck, R., Yeheskel, O., Frage, N., 2018. Addit. Manuf. 24, 257. With kind permission of Elsevier.)

401

Time-dependent deformation creep

Fig. 7.28 Creep curves of SLM AlSi10Mg specimens under various stresses at 225°C. (Uzan, N.E., Shneck, R., Yeheskel, O., Frage, N., 2018. Addit. Manuf. 24, 257. With kind permission of Elsevier.)

–8

225°C 117MPa

–9

250°C 275°C 300°C

In(Strain rate)

–10 –11 –12

-Q/R

–13

1

–14 –15 –16 17

18

19 4

10 /T, K

20

21

-1

Fig. 7.29 Creep rate as function of the reciprocal temperature at 117 MPa. (Uzan, N.E., Shneck, R., Yeheskel, O., Frage, N., 2018. Addit. Manuf. 24, 257. With kind permission of Elsevier.)

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Additive and traditionally manufactured components

–8 –9

117MPa 127MPa

225°C

137MPa 147MPa

In(Strain rate)

–10 –11 –12

n=25

–13

1

–14 –15 –16 18.5

18.6

18.7

18.8

18.9

In(Stress) Fig. 7.30 Creep rate as function of applied stress at 225°C. (Uzan, N.E., Shneck, R., Yeheskel, O., Frage, N., 2018. Addit. Manuf. 24, 257. With kind permission of Elsevier.)

The value of the stress exponent derived and indicated in Fig. 7.30 is n ¼ 25, a high value relative to the one usually observed. Even in CG Ti6Al4V the value of n is 16.7 (see Fig. 7.23) which drops to a value in the range 4–5 in UFG structures. It is likely that the creep mechanism is associated with an intergranular dislocation process controlled by dislocation glide and climb. Since climb requires diffusion and the temperature is relatively low, although influenced by the applied stress, the author believed that grain boundary diffusion might be associated with the climb process. The microstructure evolution during creep is illustrated in Fig. 7.31. The high value of the stress exponent, n, mentioned earlier is thought to be attributed to the microstructural changes as shown in Fig. 7.31a and b involving the formation and separation of submicron Si particles. The creep takes place in the Al matrix while the Si particles remain undeformed (12%) and act as a second phase strengthener leading to the high value of n. Additional structural information can be from Fig. 7.32. Total failure of material exposed to creep occurs due to creep rupture. Therefore, it is essential to evaluate the lifetime of materials exposed to creep. One of the common methods to predict extended-time

Time-dependent deformation creep

403

Fig. 7.31 Microstructure evolution during heat treatment: (a) stress-relieved specimens with continuous Si-rich cellular boundaries, (b) thin and partly broken cellular boundaries after creep at 117 MPa and 225°C, (c) formation of separated submicron Si particles after creep at 117 MPa and 300°C. (Uzan, N.E., Shneck, R., Yeheskel, O., Frage, N., 2018. Addit. Manuf. 24, 257. With kind permission of Elsevier.)

creep lives from short-time tests is the Larson-Miller [see Eq. (7.16)] approach. The relation between the rupture time and the inverse temperature and the relation between stress and the Larson parameter are presented in Fig. 7.33. It might be of interest to add additional information on creep in Al10SiMg complementary to the data presented earlier. Creep strain-time curves of horizontal samples to building direction are illustrated in Fig. 7.34 for the test conditions: (a) 180°C/200 MPa, (b) 150°C/ 200 MPa, and (c) 180°C/150 MPa. Curve (a) indicates all three stages of the creep in Al10SiMg. The alloy was prepared by SLM which is known as an AM process. The fractography of a crept sample tested to failure at 180°C/200 MPa shows cracks apparently mostly aligned in the same direction, which is illustrated in Fig. 7.35. The sample ruptured at 18.7 h. The predicted rupture time for the same alloy was 14.8 h by using the Larson-Miller plot. Fig. 7.36 shows a plot of porosity versus the energy density. Below 50 J/mm3, i.e., at low energy density, the porosity is

404

Additive and traditionally manufactured components

Fig. 7.32 Typical EBSD orientation maps of microstructure after stress relief treatment: (a) creep at 225°C, 117 MPa and (b) creep at 300°C, 117 MPa (c). (Uzan, N.E., Shneck, R., Yeheskel, O., Frage, N., 2018. Addit. Manuf. 24, 257. With kind permission of Elsevier.)

high, probably due to the lack of consolidation. It is likely that the presence of oxide layer prevented bonding. Then the amount of porosity decreases with increasing energy density. However, at 60 J mm3 and above, a large scatter is observed in the porosity values. Thus, there is a certain threshold energy density that gives maximum density for the Al alloy, which is 60–75 J/mm3. It seems from the fractured surfaces that the presence of significant amounts of unmelted powder give rise to local cracking. Clearly, increase in porosity and amount of cracking influence not only the tensile properties but also the creep resistance of materials including those of Al10SiMg.

405

Time-dependent deformation creep

a 2

225°C 250°C 275°C 300°C log10(tr) = A/T-C

1 log10(tr), h

σ = 117MPa C = 19.08

0

-1

17.0

17.5

18.0

18.5

19.0 4

10 /T, K

b

19.5

150

20.5

225°C 250°C 275°C 300°C

145 140 Stress, MPa

20.0

-1

135 130 125 120 115 110 9000

9250

9500

9750

10000

10250

10500

P=T(log10(tr)+19.08)

Fig. 7.33 (a) Relationship between creep rupture time (tr) and reciprocal temperature under an applied stress of 117 MPa and (b) relationship between experimental stress and the Larson-Miller parameter at rupture. (Uzan, N.E., Shneck, R., Yeheskel, O., Frage, N., 2018. Addit. Manuf. 24, 257. With kind permission of Elsevier.)

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Additive and traditionally manufactured components

4.5 4

(a)

Strain/%

3.5 3 2.5 2 1.5 (b)

1

(c)

0.5 0 0.00

5.00

10.00

15.00

20.00

Time/h

Fig. 7.34 Creep curves of SLM fabricated AlSi10Mg alloy (horizontal samples) at the following conditions: (a) 180°C/200 MPa, (b) 150°C/200 MPa, and (c) 180°C/150 MPa. (Read, N., Wang, W., Essa, K., Attallah, M.M., 2015. Mater. Des. 65, 417. With kind permission of Elsevier.)

Fig. 7.35 Backscattered SEM fractographs of SLM-fabricated AlSi10Mg horizontal samples, showing sample tested at 180°C/200 MPa. (Read, N., Wang, W., Essa, K., Attallah, M.M., 2015. Mater. Des. 65, 417. With kind permission of Elsevier.)

Time-dependent deformation creep

407

Fig. 7.36 Porosity variation vs the energy density. (Read, N., Wang, W., Essa, K., Attallah, M.M., 2015. Mater. Des. 65, 417. With kind permission of Elsevier.)

7.7 Tensile creep in CP Al AA6061 Despite a large body of literature and the efforts devoted to understand the creep behavior of aluminum alloys, in additive manufactured AA6061 (a most popular Al alloy), a full description of this phenomenon on the basis of microstructural parameters and experimental mechanical behavior is, currently, still missing. Therefore, tensile creep properties of the conventionally manufactured AA6061 are presented in the following. Creep tests on Al 6061 (T6) were performed at the stresses of 65.4 and 98.1 MPa at temperatures of 573, 623, and 673 K. Constitutive model was applied to predict and compare the creep behavior using experimental material constants for the creep constitutive equations. The determination of the material constants within the creep constitutive equations are discussed in the following. Comparison of the analytical and experimental creep curves for the stress-temperature combinations of 65.4 MPa-673 K, 98.1 MPa573 K, 98.1 MPa-623 K, and 98.1 MPa-673 K, respectively, is illustrated in Figs. 7.37–7.40. Note that in these figures all the three creep stages are present, thus the tertiary creep (accelerated creep stage) observed leads to the creep rupture of the Al 6061 samples. Several parameters are of interest. Thus Eq. (7.18) can be used to evaluate the stress exponent, n. The slope of the plot on a log-log scale provides the value of n. The log-log expression of Eq. (7.18) is lnðε_ Þ ¼ ln A + ln σ + nðQ=RT Þ

(7.20)

408

Additive and traditionally manufactured components

0.18 Experimental creep curve, for applied stress 65.4 MPa at temperature 673 K

0.16 0.14

Analytical creep curve, for applied stress 65.4 MPa at temperature 673 K

Creep strain

0.12 0.1 0.08 0.06 0.04 0.02 0 -0.02

0

50

100

150

200

250

300

350

400

Time, s Fig. 7.37 Comparison of experimental and analytical creep curves for applied stress 65.4 MPa and at temperature 673 K (400°C). (Shivakumar, S.P., Sharan, A.S., Sadashivappa, K., 2018. IOP Conf. Ser.: Mater. Sci. Eng. 310, 012027; Open access.)

0.08 0.07

Creep strain

0.06 0.05 0.04 0.03 Experimental creep curve, for applied stress 98.1 MPa at temperature 573 K

0.02

Analytical creep curve, for applied stress 98.1 MPa at temperature 573 K

0.01 0 0

10000

20000

30000

40000

50000

60000

70000

80000

90000

Time, s

Fig. 7.38 Comparison of experimental and analytical creep curves for applied stress 98.1 MPa and at temperature 573 K (300°C). (Shivakumar, S.P., Sharan, A.S., Sadashivappa, K., 2018. IOP Conf. Ser.: Mater. Sci. Eng. 310, 012027; Open access.)

A plot of logarithms of the strain rate against the logarithm of applied stress is illustrated for Al 6061. The value of n varies in the range 1–1.33 depending on the combination of the applied stress and temperatures. The plot of Fig. 7.41 is for 65.4 MPa stress and that of Fig. 7.42 is for 98.1 MPa.

409

Time-dependent deformation creep

0.18 0.16 0.14

Creep strain

0.12 0.1 0.08 0.06 Experimental creep curve, for applied stress 98.1 MPa at temperature 623 K

0.04 0.02

Analycal creep curve, for applied stress 98.1 MPa at temperature 623 K

0 -0.02 0

500

1500

1000

2000

2500

Time, s

Fig. 7.39 Comparison of experimental and analytical creep curves for applied stress 98.1 MPa and at temperature 623 K (350°C). (Shivakumar, S.P., Sharan, A.S., Sadashivappa, K., 2018. IOP Conf. Ser.: Mater. Sci. Eng. 310, 012027; Open access.)

0.14 0.12

Creep strain

0.1 0.08 0.06 0.04

Experimental creep curve, for applied stress 98.1 MPa at temperature 673 K

0.02

Analytical creep curve, for applied stress 98.1 MPa at temperature 673 K

0 0

50

100

150

200

250

300

-0.02

Time, s

Fig. 7.40 Comparison of experimental and analytical creep curves for applied stress 98.1 MPa and at temperature 673 K (400°C). (Shivakumar, S.P., Sharan, A.S., Sadashivappa, K., 2018. IOP Conf. Ser.: Mater. Sci. Eng. 310, 012027; Open access.)

From the slope of ln (instantaneous strain rate) versus instantaneous stress, the value of the material constant β can be calculated. The plots are shown in Figs. 7.43 and 7.44 for the indicated temperatures at stress of 64.5 and 98.1 MPa, respectively. From these plots, it can be observed that the derived values of the material parameter β are in the range

410

Additive and traditionally manufactured components

0

In(Instantaneous strain rate), S-1

For applied stress 65.4 MPa at temperature 573 K -2 For applied stress 65.4 MPa at temperature 623 K -4 For applied stress 65.4 MPa at temperature 673 K -6 -8 -10 -12 -14 -16 0

500

1000

1500

2000

In(Instantaneous stress), MPa Fig. 7.41 Plots of ln (instantaneous strain rate) (s1) against ln (instantaneous stress) (MPa), for applied stress 65.4 MPa and at varying temperature 573 K (300°C), 623 K (350°C), and 673 K (400°C). (Shivakumar, S.P., Sharan, A.S., Sadashivappa, K., 2018. IOP Conf. Ser.: Mater. Sci. Eng. 310, 012027; Open access.)

In(Instantaneous strain rate), S-1

0 For applied stress 98.1 MPa at temperature 573 K -2

For applied stress 98.1 MPa at temperature 623 K

-4

For applied stress 98.1 MPa at temperature 673 K

-6 -8 -10 -12 -14 0

2

4

6

8

10

In(Instantaneous stress), MPa Fig. 7.42 Plots of ln (instantaneous strain rate) (s1) against ln (instantaneous stress) (MPa) for applied stress 98.1 MPa and at varying temperature 573 K (300°C), 623 K (350°C). (Shivakumar, S.P., Sharan, A.S., Sadashivappa, K., 2018. IOP Conf. Ser.: Mater. Sci. Eng. 310, 012027; Open access.)

β ¼ 3.2068  103–6.80  103 for various combination of the applied stress and the temperature. The value of α can be calculated from the relation. α¼

β n

(7.21)

411

Time-dependent deformation creep

0

In(Instantaneous strain rate), S-1

For applied stress 65.4 MPa at temperature 573 K -2 For applied stress 65.4 MPa at temperature 623 K -4 For applied stress 65.4 MPa at temperature 673 K -6 -8 -10 -12 -14 -16 0

500

1000

1500

2000

Instantaneous stress, MPa

In(Instantaneous strain rate), s-1

Fig. 7.43 Plots of ln (instantaneous strain rate) (s1) against instantaneous stress (MPa), for applied stress 65.4 MPa and at varying temperature 573 K (300°C), 623 K (350°C) and 673 K (400°C). (Shivakumar, S.P., Sharan, A.S., Sadashivappa, K., 2018. IOP Conf. Ser.: Mater. Sci. Eng. 310, 012027; Open access.)

0 For applied stress 98.1 MPa at temperature 573 K -2

For applied stress 98.1 MPa at temperature 623 K

-4

For applied stress 98.1 MPa at temperature 673 K

-6 -8 -10 -12 -14

0

500

1000

1500

2000

Instantaneous stress, MPa Fig. 7.44 Plots of ln (instantaneous strain rate) (s1) against instantaneous stress (MPa) for applied stress 98.1 MPa and at varying temperature 573 K (300°C), 623 K (350°C), and 673 K (400°C). (Shivakumar, S.P., Sharan, A.S., Sadashivappa, K., 2018. IOP Conf. Ser.: Mater. Sci. Eng. 310, 012027; Open access.)

The Larson-Miller parameter (Eq. 7.16) can be used to calculate the activation energy, Q, for creep. The value of Q was found to vary in the range 165.39  103 J/mol to 177.98  103 J/mol. It would be appropriate to

412

Additive and traditionally manufactured components

derive the Larson-Miller parameter starting with relation of (7.18) (Sherby and Dorn, 1952) rewritten here as   Q n ε_ ¼ Aσ exp  (7.22) RT The rupture strain is obtained by integrating Eq. (7.18)   Q n1 t ¼ tr ¼ constant σ exp  RT

(7.23)

The time, t, is equated to the creep rupture time, tr. Now taking logarithm on both sides of Eq. (7.23) results in log ðtr Þ ¼ log ðconstantÞ  n log ðσ Þ +

Q 1 2:3R T

(7.24)

Arrange Eq. (7.24) to obtain T log ðtr Þ ¼ T ½ log ðconstantÞ  n log ðσ Þ +

Q 2:3R

(7.25)

By a constant stress value, Eq. (7.25) can be written as T log ðtr Þ  T log ðc Þ ¼

Q 2:3R

(7.26)

and rearranging Eq. (7.26), with log(c) being a constant, C, results in Q ¼ PLM ¼ T ½ log ðtr Þ + C Þ 2:3R

(7.27)

where T is given in hours, T in the absolute temperature. Usually, C is 20, but for Al a value of 14 is suggested. It might be of interest to add creep data on an earlier work reviewing the mechanical properties of Al alloys. The creep behavior of Al 6061 T651 (recall that the designation of T651 refers to solutionized, stress-relievedstretched and artificially aged alloy) is shown in Fig. 7.45 at various temperatures. The primary creep rate stage in this figure is insignificant and the second-stage creep dominates the curves. An often used relation suggested by Garofalo is used for creep behavior experiments. It is a hyperbolic sine function in the form of   Q ε_II ¼ A exp  (7.28) ½ sinh ðBσ Þn RT

413

Time-dependent deformation creep

0.6 0.5

0.010 0.005 0.000

0.8

0.020 150 MPa 140 MPa 130 MPa

0.7

0.015

0

2000 4000 Time (s)

0.4

6000

0.3 0.2 0.1

0.6 0.5

0.015 0.010 0.005 0.000 0

1000 2000 3000 4000

0.4

Time (s)

0.3 0.2 0.1

0.0 0

2000

0.0

6000

4000

0

1000

2000

Time (s)

0.5

4000

5000

0.020 50 MPa 45 MPa 40 MPa

0.7

0.015 0.010 0.005 0.000 0

2000 4000 6000 8000 Time (s)

0.4 0.3 0.2

Creep Strain

Creep Strain

0.6

0.8

d

0.020 80 MPa 70 MPa 60 MPa

0.7

3000

Time (s)

Creep Strain (mm/mm)

C 0.8 Creep Strain (mm/mm)

Creep Strain

Creep Strain

Creep Strain (mm/mm)

b

0.020 220 MPa 210 MPa 200 MPa

0.7

Creep Strain (mm/mm)

a 0.8

0.6 0.5

0.015 0.010 0.005 0.000 0

500 1000 1500 2000 2500

Time (s)

0.4 0.3 0.2

0.1

0.1

0.0

0.0 0

2000

4000

6000

8000

10000

12000

0

1000

Time (s)

e 1.0

3000

4000

0.020 20 MPa 15 MPa 13 MPa

0.8

Creep Strain

Creep Strain (mm/mm)

2000 Time (s)

0.6

0.015 0.010 0.005 0.000 0

500

1000 1500 2000 Time (s)

0.4

0.2

0.0 0

1000

2000

3000

4000

5000

Time (s)

Fig. 7.45 Creep behavior of 6061-T651 at (a) 200°C, (b) 250°C, (c) 300°C, (d) 350°C, and (e) 400°C. The inset figures detail the creep behavior at strains