Lightweight Alloys for Aerospace Applications E-Book

• Herausgeber: John Wiley & Sons
• Kategorie: Fachliteratur
• Sprache: Englisch

Seitenzahl: 454

Veröffentlichungsjahr: 2013

Beschreibung

This proceedings volume from the 2001 TMS Annual Meeting & Exhibition covers advances made in the area of scientific understanding of technological application of lightweight alloys. Papers focus on fundamental science as well as application and concentrate on scientific advances in aluminum, magnesium, titanium, and beryllium alloys and their composites. Processing, structure-property relationship, failure mechanisms, and advanced joining themes are also discussed.

Leseprobe

Contents

Cover

Half Title page

Title page

Copyright page

Foreword

Acknowledgements

Aluminum Alloys

Precipitation Hardening - The Oldest Nanotechnology

Abstract

The story of a discovery

Lattice correspondences

Metastable phases

Extrinsic defects

Rules for the creation of nano-size dispersoids

Hardening

References

The Role of Ledge Nucleation/Migration in Ω Plate Thickening Behaviour in Al-Cu-Mg-Ag Alloys

Abstract

Acknowledgments

1. Introduction

2. Experimental Procedure

3. Results

4. Discussion

5. Conclusions

6. References

Improving Recrystallization Resistance in Wrought Aluminum Alloys with Scandium Addition

Abstract

Introduction

Experimental

Results

Discussion

Conclusion

References

Application of 3D Digital Image Processing to Quantify Fracture Micro-Mechanisms in Al 7050 Alloy

Abstract

Introduction

Background

Experimental Work

Conclusions

Acknowledgements

References

On the Effect of Stress on Nucleation, Growth, and Coarsening of Precipitates in Age-Hardenable Aluminum Alloys

Abstract

Introduction

Experimental Procedure

Results

Discussion

Acknowledgement

References

Hot Deformation Mode and TMP in Aluminum Alloys

Introduction

Hot Extrusion

Hot Rolling

Forging

Conclusion

References

Microstructural Characterization of Friction Stir Welds

Abstract

Introduction

Experimental Procedure

Results

Discussion

Conclusions

References

Acknowledgements

Deformation, Fracture and Fatigue in A Dispersion Strengthened Aluminum Alloy

Abstract

Introduction

Results

Discussion

Summary and Conclusion

References

The Effect of Retrogression and Reaging on 7249 Aluminum Alloy

Abstract

Introduction

Experimental Procedure

Results and Discussion

Conclusions and Recommendations

Acknowledgements

References

Aluminum-Lithium Alloys

Grain Boundary Corrosion and Stress Corrosion Cracking Studies of Al-Li-Cu Alloy AF/C458

Abstract

Introduction

Experimental Procedures

Results and Discussion

Summary

Acknowledgements

References

Localized Corrosion Mechanisms of Al-Li-Cu Alloy AF/C458 After Interrupted Quenching From Solutionizing Temperatures

Abstract

Introduction

Experimental Procedures

Results and Discussion

Summary

Acknowledgments

References

Effect of Single and Duplex Aging on Microstructure and Fatigue Crack Growth in Al-Li-Cu Alloy AF/C-458

Abstract

Introduction

Background

Experimental

Results

Microstructure of sample aged at 150°C for 3 hours:

Discussion

Conclusions

Acknowledgements

References

Effect of Friction Stir Welding on the Superplastic Behavior of Weldalite Alloys

Abstract

Introduction

Materials and Experimental Procedure

Results

Discussion

Conclusion

References

On the Effect of Thermal Exposure on the Mechanical Properties of 2297 Plates

Abstract

Introduction

Experimental Procedure

Experimental Results

Conclusions

Acknowledgement

Reference

Titanium Alloys

The Effect of Crystal Orientation and Boundary Misorientation on Tensile Cavitation During Hot Tension Deformation of Ti-6Al-4V

Abstract

Introduction

Experimental Procedures

Results and Discussion

Conclusions

Acknowledgements

References

Processing and Properties of Gamma Titanium Aluminides and Their Potential for Aerospace Applications

Abstract

Introduction

Alloy development

Ingot manufacture

Thermomechanical processing

Joining and Machining Techniques

Component fabrication

Summary

Acknowledgements

References

Metallurgical and Fabrication Factors Relevant to Fracture Mechanism of Titanium –Aluminide Intermetallic Alloys

Abstract

Introduction

Conclusions

References

Dwell-Fatigue Behavior of Ti-6Al-2Sn-4Zr-2Mo-0.1Si Alloy

Abstract

1. Introduction

2. Material and Microstructures

3. Experimental Procedures

4. Results and Discussion

5. Summary and Conclusions

Acknowledgments

References

Characterization of Microstructural Evolution in a Ti-6Al-4V Friction Stir Weld

Abstract

Introduction

Experimental Procedure

Results

Discussion

Conclusions

References

Acknowledgements

Microstructural Evolution During Hot Working of Ti-6Al-4V at High Strain Rates

Abstract

Introduction

Experimental

Results and Discussion

Conclusions

Acknowledgements

References

The Application of a Novel Technique to Examine Sub-β Transus Isothermal Forging of Titanium Alloys

Abstract

Introduction

Experimental Procedure

Results & Discussion

Compression Flow Curves

Conclusions

Acknowledgements

References

Stress Corrosion Cracking of α-Ti in a Methanol Solution

Abstract

Introduction

Experimental Procedure

Experimental Results

Discussion

Summary

Acknowledgments

References

Figure Caption

Investment Casting of Titanium Alloys with CaO Crucible and CaZrO3 Mold

Abstract

Introduction

Experimental procedure

Results and discussion

Conclusion

References

Composites

Compressive Behavior of Ti-6Al-4V/TiC Layered Composites: Experiments and Modeling

Abstract

Introduction

Experimental Procedure

Experimental Results

Model

Summary and Conclusions

Acknowledgments

References

High Ductility Cast Aluminum Beryllium Alloys

Abstract

Introduction

Experimental Procedures

Results and Discussion

Summary

References

Characterization of Reinforcing Particle Size Distribution in a Friction Stir Welded Al-SiC Extrusion

Abstract

Introduction

Material System

Experimental procedures

Statistical Characterization

Summary and Conclusions

Acknowledgements

References

In-Situ Formation of AIN Reinforced Al Alloy Composites Using Ammonia

Abstract

Introduction

Thermodynamic Modeling

Experimental

Results and Discussion

Conclusions

Acknowledgements

References

Author Index

Subject Index

LIGHTWEIGHT ALLOYS FORAEROSPACE APPLICATION

Edited by:Dr. Kumar Jata, Dr. Eui Whee Lee, Dr. William Frazier and Dr. Nack J. Kim

Partial funding for this publication was provided by the Seeley W. Mudd Fund.

A Publication of The Minerals, Metals & Materials Society 184 Thorn Hill Road Warrendale, Pennsylvania 15086-7528 [phone]

Visit the TMS web site athttp://www.tms.org

The Minerals, Metals & Materials Society is not responsible for statements or opinions and is absolved of liability due to misuse of information contained in this publication.

ISBN Number 0-87339-491-7Library of Congress Number: 2002108586

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© 2001

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FOREWORD

This manuscript contains a collection of 27 papers presented at the Lightweight Alloys for Aerospace Applications symposium at the TMS Annual Meeting in New Orleans, February 12-15, 2001. The manuscript contains outstanding papers on:

Current understanding of coarsening of precipitates in novel aluminum alloys

Deformation, fracture and fatigue and corrosion resistance of aluminum-lithium alloys

Nano-crystalline aluminum alloys

Processing, fatigue and fracture and relationships to microtexture in Ti

Status of gamma titanium aluminides

Microstructure and properties of metal matrix composites

ACKNOWLEDGEMENTS

The organizers of the symposium would like to extend their sincere appreciation to Professors Erhard Hornborgen, Gary Shiflet, Thomas Sanders, Hugh McQueen, Arun Gokhale and to Drs. Jonathan Paul, Kamran Nikbin and Awadh Pandey for participating in the symposium as invited speakers and to Dr. Mary C. Juhas and Prof. O.S. Es-Said who helped in chairing the sessions. The organizers would like to thank all the participants who made this symposium a big success.

Kumar V. Jata Air Force Research Laboratory, USA June 2001

LIGHTWEIGHT ALLOYS FOR AEROSPACE APPLICATION

Edited by:Dr. Kumar Jata, Dr. Eui Whee Lee,Dr. William Frazier and Dr. Nack J. Kim

ALUMINUM ALLOYS

Precipitation Hardening – The Oldest Nanotechnology

Erhard Hornbogen

Pgs. 1–11

184 Thorn Hill Road Warrendale, PA 15086-7514 [phone]

PRECIPITATION HARDENING - THE OLDEST NANOTECHNOLOGY

Erhard Hornbogen

Ruhr-University Bochum, Institute for Materials 44780 Bochum, Germany

Abstract

Precipitation hardening of aluminum was discovered about 100 years ago by Dr. Alfred Wilm [1,2]. Using aluminum alloys as example a survey is given on mechanism and limits of precipitation hardening. It is discussed how hard, nanometer-size particles (nanos, greek, the dwarf) can form as an ultra fine dispersoid. A simple example for optimum conditions is provided by diamond cubic particles (Si, Ge) in the f.c.c. Al-Matrix. The role of a sequence from more to less metastable phases is discussed, as well as the effects of additional (and trace) alloying elements. The amount of precipitation hardening is limited, besides by the volume fraction of particles, by their strength. This in turn determines the critical diameter above which the transition from passing to by-passing by dislocations takes place. Simple models for the calculation of this microstructural parameter are discussed.

From combinations of precipitation hardening with other hardening mechanisms the limits for ultra high strengths are defined.

The story of a discovery

It is now about 100 years ago that Dr. Alfred Wilm started experiments with a wide range of aluminum alloys at the metallurgical department of the Central Institute for Scientific and Technological Studies in Neubabelsberg (close to Berlin). Major motivation for his work was to increase the insufficient strength of this then still relatively young material. By 1906 he had developed a new type of alloy with 3.5-5.5% Cu and less than 1% Mg and Mn with a strength of more than 400 MPa. It was soon well known by its trade name Duralumin. The prefix contains a possibly intended ambiguity: durus (latin, hard), but also Dürener Metallwerke, (Rhineland) the industrial firm, where the alloy was produced and shaped. [1,2]

There are the options for evolutionary, predictable progress in science and technology or revolutionary developments. The discovery of precipitation hardening was unpredicted.

Alfred Wilm knew all the physical metallurgy of his days. He knew that metals could be solid solution hardened and work hardened. He also knew that the carbon steels had to be quenched for hardening. Therefore he melted a large number of alloys which he investigated in an as-quenched, slowly cooled and worked state. To his great frustration he found out that – different from steels – some alloys became even softer by quenching.

However, one day, many quenching experiments were performed on Saturday morning. Not all the hardness measurements were completed the same day. The sun was shining and Wilm went out for the week-end to go sailing. Next monday morning the hardness measurements were completed, and surprisingly hardness as well as all strength properties had increased considerably. At first, the technician was blamed for sloppy work, but then all the measurements were repeated carefully. In addition, duration and temperature of aging was varied systematically.

1. Wilm had to follow the wrong hypothesis of the analogy between steel and aluminum to expect hardening by rapid quenching.

2. He had to be lazy, not to complete his measurements immediately, to realize the importance of aging.

3. He had to recognize the completely unexpected result, reproduce, optimize and apply it.

He could, however, not understand that the production of a fine dispersoid of nm-size particles is the physical cause of precipitation hardening. So he did not know that he discovered the first nano-technology.

Lattice correspondences

Considerable strength of Al-alloys is always due to precipitation hardening. This requires the formation of an even and ultra-fine dispersoid of hard particles, The optimum would be achieved by hard diamond cubic particles in the fcc Al-based matrix. Size and spacing should be in the nm-range. In fact, often either a fine microstructure of soft particles emerges or a coarser one of strong particles. Both do not lead to the desired hardening effect. (Fig. 1)

Figure 1: Anomalies of different physical properties due to nm-size particles

This paper will discuss the mechanisms of precipitation of second phases, using fcc solid solutions of Al as examples. They have been well explored in the past 50 years [3, 4]. The results can be interpreted by a generalized Ostwald step rule. In addition dislocation theory provides the principles on how to obtain optimum strengthening.

Finally a comparison must be made with respect to coherency between the structure of the matrix crystal α and the newly formed phases βi. This relationship may imply:

a) non-coherency, b) partial coherency, c) constrained coherency, d) full coherency (Figs. 2, 3). These possibilities provide a decreasing amount of the structural term of the interfacial energy γαβ, which essentially controls the nucleation behavior of βi from α [5, 6, 7] (Table 1). The crystal structures of strong particles (termodynamically stable and high resistance to dislocation motion: Θ−Al2Cu, δ-AlLi, Si) do not fulfill the prerequisites for coherency with fcc Al (Fig. 3).

Figure 2: Types of coherency between matrix α and particle β, (a) coherent, ordered, (b) constrained coherent (shear), (c) partially coherent, (d) non coherent

Figure 3: Formation of non-coherent Si from fcc Al (Si) solid solution

Table I Nucleation mechanisms in Al-Cu- alloys

Metastable phases

The sequence of the precipitate phases βi for the Al-Cu-alloys implies a decreasing coherency with increasing thermodynamic al stability (Fig. 4): Θ” → Θ’ → Θ.

Figure 4(a): Sequence of metastable, coherent → non-coherent Θ−Al2Cu in AlCu-alloys

Figure 4(b): Effect of Mg, Ag, on the transition to Θ−AlCu from metastable S-Al2, Mg, Ag

At a certain amount of undercooling the relatively smaller driving force becomes sufficient for a quicker formation of the metastable phase, relative to the more stable one. Its nucleation is favoured by the formation of a low interfacial energy γαβ with GN activation energy of nucleation, gβ specific free energy change Jm−3, i number of atoms in the nucleus, [6]:

(1a)

(1b)

(1c)

Thus crystallographic coherency and classical nucleation theory explain why the less stable phase forms first and not the one which leads to the most stable state. There are two reasons why multi-stage reactions cannot be expected: 1. No metastable phases exist (Al-Si, Al-Ge, Al-Zn) [8], 2. The most stable phase can form coherently. Examples for the second case (as (γ + γ)-Ni alloys) do not exist for Al-alloys.

Extrinsic defects

This in turn leads to particular combinations of lattice defects and interfaces [5, 6, 7, 8], which provide minimum activation barriers and consequently maximum rates of formation (Table 1). Such favourable nucleation processes occur only at sites where lattice defects are preexisting. Consequently the bulk rate depends on grain size or dislocation density. This leads to the phenomenon that different phases form simultaneously, but at the various sites inside the alloy. During particle growth small particles dissolve in favour of large ones. But even more important, less stable phases are dissolved in the environment of more stable phases due to the differences in local solubilities. Consequently particle-free zones and uneven distribution of particles in the interior of grains develop (Fig. 5).

Figure 5: Defect-induced nucleation of different phases (compare Table 1), formation of particle free zones.

At any defect:

(2a)

At dislocation:

(2b)

At boundary:

(2c)

Besides lattice defects some trace elements are known to have favorable effects on the dispersoid microstructure. Wilm reported already the effect of Mg on the binary Al-Cu-alloy. An additional effect of Ag was found later on [4]. Such elements should favor the formation of a finer dispersoid. This can be explained by solubility of the trace element in the coherent metastable phase, while the stable non-coherent phase shows no solubility [9]. This in turn retards anomalous coarsening by the in-situ-transformation, coherent → non-coherent.

Rules for the creation of nano-size dispersoids

Unconstrained coherent nucleation comes closest to homogeneous nucleation and therefore to the formation of the finest possible dispersoid of the second phase βi [10]. However, a dispersoid of a incoherent, hard phase is desired.

For Al-alloys there is usually a multitude of options for reactions which lead closer to equilibrium at a wide range of rates (Equ. 1, 2). For Al-Cu the following possibilities exist: The stable phase Θ and three metastable phases, combined with three types of lattice defects and eventually a continuous and a discontinuous mode; i.e. there is a competition between 13 options. The fastest ones will win and produce the microstructure. The thermodynamical principle behind this phenomenon seems to be the maximization the initial rate of entropy production. The eventual creation of structural order (−Sstr) is overcompensated by the fast production of thermal entropy +Sth. This in turn is a generalisation of Ostwalds step rule [10].

(1.3)

The initial, microstructure can be interpreted in terms of Equ. 3. This relation is favored by the following set of sometimes conflicting properties:

* this is favored in ternary alloys by the ±-effect, i.e. one solute larger, the other smaller than Al (Table 2).

Table II Atomic size ratio of solutes in Al

Smaller

Larger

Cu

Ti

Zn

Sn

Si

Cd

Ni

Ge

Co

Li

Fe

Mg

Mn

Ga

Cr

Hardening

The formation of a fine dispersoid of particles is not sufficient for hardening [11,12]. Necessary for the validity of the OROWAN-equation (Figs. 6, 7, 8) are strong particles which sustain the stress excerted by moving dislocations, d > dC:

Figure 6: Calculated (Equ. 4) upper limit of hardening by different volume fractions of particles in Al.

Figure 7: Definition of size ranges and critical particle diameters dC, for 1% particles

Figure 8: Geometrical features and particle-dislocation interactions for dispersoid microstructures

(4)

there is always a critical particle diameter dC below which the particles are sheared. This diameter should be as small as possible (Table 3). The critical diameter dC defines the maximum hardening effect by a certain volume fraction of particles (Fig. 6). It can be estimated by examining the force F which a single particle is able to excert on a looping dislocation. For F < Gb2, shearing occurs and by-passing for F ≥ Gb2 (Figs. 7, 8).

(5a)

and C≈1 is a geometrical factor depending on the shape of the particle. G and b are shear modulus and Burgers vector of the fcc-matrix solid solution.

For disordered coherent particles no APB has to be created, but the amount of the difference in critical shear stress between matrix α and particle β (τα-τβ) becomes relevant for modest hardening:

(5b)

Pores or liquid inclusions are sheared at any size, inspite of their strong hardening effect (Equ. 4):

(5c)

Particles β with a different crystal structure from the matrix α require the theoretical shear stress to create a dislocation bβ. Consequently only very small incoherent particles are sheared:

(5d)

Low dC-particles like Si will provide the highest hardening pro volume fraction.

These equations have to be modified if not one but a pair or more dislocations interact with a particle. It follows that the “art” of causing precipitation hardening implies the production of even and (S → min, Equ. 4) fine dispersoids of particles with small critical sizes dC (Table 3).

Usually, it is easy to form small subcritical particles d < dC. As they are sheared they must cause less hardening than the by-passing mechanism. Also consequences on localization of strain have to be considered, for example on initiation and propagation of cracks under fatigue or stress corrosion conditions.

Finally it has to be considered that ultra high strength cannot be obtained by precipitation hardening Δσp alone. A high yield stress σy must be built up from contributions of additional hardening mechanisms. They can be systematically discussed by considering the 0- to 2-dimension of obstacles to the motion of dislocations [12].

(6)

Where σ0 is the very low strength of pure Al, and Δσs the contribution of solid solution hardening, Δσd of a dislocation forest, Δσb of grain boundaries. The different terms are not independent of each other. For our discussion it is important that for d > dC, fine grain hardening Δσb becomes irrelevant. Δσs and Δσd are used to built up high strength of Al-alloys. However, precipitation hardening Δσp contributes usually the biggest share to high strength. This in turn is always due to strong particles in the size range between 1 and 10 nano-meters (Table 3).

Table III Examples for critical particle diameters dc in Al

References

1. A. Wilm, Metallurgie, 8 (1911), 223

2. C. Kammer, “Success for aluminium thanks to 75 years of materials research”, Aluminium, 75 (1999) 753–775

3. I. W. Martin, Micromechanisms in Particle-Hardened Alloys (Cambridge, UK: Cambridge Univ. Press, 1979)

4. I. J. Polmear, Light Alloys, 3rd ed. (London, UK: Edward Arnold, 1995), 41, 105

5. E. Hornbogen, “Nucleation of precipitates in defect solid solutions”, Nucleation, ed. A.C. Zettlemoyer (New York, NY: Marcel Dekker, 1969), 309–378

6. E. Hornbogen, “Electronmicroscopy of Precipitation in Al-Cu Solid Solutions”, Aluminium, 43 (1967), 41, 115, 163, 170

7. E. Hornbogen, “Combined Reactions”, Met. Trans., 10a (1979), 947–971

8. E. Hornbogen, A. K. Mukhopadhyay and E.A. Starke, “An Exploratory Study of Hardening of Al-(Si, Ge)–alloys”, Z. Metallkunde, 82 (1992), 577–580

9. A. K. Mukhopadhyay, “Compositional Characterization of Cu-rich Phase Particles Present in As-Cast Al-Cu-Mg-alloys Containing Ag”, Mat. Trans. A, 30A (1999), 1693–1704

10. E. Hornbogen, “Formation of nm-size dispersoids from Supersaturated Solid Solutions of Al”, Materials Science Forum, Vol. 331–337 (2000), 879–888

11. C. P. Blankenship, E. Starke and E. Hornbogen, “Microstructures and properties of Al-alloys”, Microstructure and properties of materials 1, ed. J.C.M. Li (Singapore: World Scientific 1996) 1–51

12. E. Hornbogen, E. Starke, “High Strength low alloy aluminum”, Acta Mat., 41 (1993), 1–16

LIGHTWEIGHT ALLOYS FOR AEROSPACE APPLICATION

Edited by: Dr. Kumar Jata, Dr. Eui Whee Lee,Dr. William Frazier and Dr. Nack J. Kim

ALUMINUM ALLOYS

The Role of Ledge Nucleation/Migration in Ω Plate Thickening Behaviour in Al-Cu-Mg-Ag Alloys

C.R. Hutchinson, X. Fan, S.J. Pennycook and G.J. Shiflet

Pgs. 13-23

184 Thorn Hill Road Warrendale, PA 15086-7514 [phone]

THE ROLE OF LEDGE NUCLEATION/MIGRATION IN Ω PLATE THICKENING BEHAVIOUR IN Al-Cu-Mg-Ag ALLOYS

1C. R. Hutchinson, 2, 3X. Fan, 3S. J. Pennycook, 1G. J. Shiflet

1Dept. of Mat. Sci. and Eng., University of Virginia, Charlottesville, VA, 22903, USA.