Development of Low-cost Casting Titanium Alloys: An Integrated Computational
Materials Engineering (ICME) Guided Study
Dissertation
Presented in Partial Fulfillment of the Requirements for the Degree Doctor of Philosophy
in the Graduate School of The Ohio State University
By
Zhi Liang
Graduate Program in Materials Science and Engineering
The Ohio State University
2018
Dissertation Committee
Professor Alan A. Luo, Advisor
Professor Glenn Daehn
Professor Ji-cheng Zhao
Copyrighted by
Zhi Liang
2018
Abstract
Titanium alloys have proved to be important lightweight structural materials since
1960’s, due to their excellent intermediate temperature mechanical properties, corrosion resistance and weldability. Their good property-to-weight ratios make them ideal for many high-end and weight-sensitive applications. However, the application of titanium alloys is still limited due to the high costs in raw materials and manufacturing, indicating the importance of developing new cost-effective titanium alloys. Compared with other lightweight structural alloys (e.g. aluminum, magnesium), the raw material cost for titanium alloys is generally considered as expensive due to its expensive alloying elements such as vanadium, molybdenum, and tin. The cost issue is further amplified by the difficulties in using conventional machining methods for component-shaping due to the low thermal conductivity. Therefore, cost-effective titanium alloys should address either aspect. This work focuses on the goal of developing new cost-effective Ti-Al-Fe-
Mn titanium alloys for the casting process via Integrated Computational Materials
Engineering (ICME) approach by using cheaper alloying elements and net-shape manufacture process.
Calculation of Phase Diagram (CALPHAD) work on the Ti-Al-Mn ternary system was conducted to establish the reliable thermodynamic database to guide the alloy design.
ii
This investigation includes both experimental work and database programming.
Isothermal phase equilibria and differential scanning calorimetry (DSC) experiments were conducted to acquire the equilibrium information for the determination of phase boundaries. Based on the experimental results, an updated Ti-Al-Mn ternary thermodynamic database was built, which is beneficial for further multi-component titanium alloy development.
The other involved ternary titanium alloy system, Ti-Al-Fe ternary system, was investigated with CALPHAD method first to determine target compositions and desired phases, and then characterized experimentally with scanning electron microscopy (SEM) in order to establish the relationship between process parameters and microstructure.
Based on the combining results from CALPHAD and selected experiments, a new Ti-Al-
Fe casting titanium alloy was designed and produced with induction melting.
Microstructure characterization and mechanical property testing were conducted to examine the potent of this new alloy and establish the preliminary structure-property relationship in the as-cast condition. Besides regular scanning electron microscopy
(SEM), transmission electron microscopy (TEM) was also applied to investigate the effect of certain nano-size microstructural features linked to mechanical properties.
Finally, a laboratory-scale manufacture framework was constructed in OSU facility, including vacuum induction skull melting (ISM), gravity tilt-pour casting, and permanent metallic mold casting. This new framework was used to manufacture a prototype casting connecting rod with the new Ti-Al-Fe casting alloy served as iii
conceptual validation for further industrialization. A preliminary cost analysis was also
presented in this work to illustrate its commercialization potent compared with the current industry applications (e.g. alloy and process).
iv
Dedication
Dedicated to my parents, Fei Huang and Liming Liang, for their continuous and patient
support.
v
Acknowledgments
I shall firstly acknowledge my advisor, Prof. Alan Luo for his full support and
guidance on my PhD research period. With his significant industry experience and
courage, we successfully stepped into a brand new field, titanium alloy, for both of us.
Being pushed into this new and challenging area of research, he offered me a precious experience of handling and planning an entire research project. My achievements in OSU and my new career in NIST are mostly thanked to his backing and recommendations.
I also want to acknowledge the constant consultation from the other co-PI/advisor during my PhD research/project, Prof. James C. Williams. It is my honor to receive his suggestions and counseling in the titanium field, which is new for both me and Prof. Luo.
With his comprehensive experience of titanium in both academia and industry, he gave us many important hints in planning and conducting experiments, and giving us revisions and feedbacks for our publications.
I would like to thank Prof. Glenn Daehn and Prof. Ji-Cheng Zhao of OSU for their help and guidance for the completion of my degree, especially making time for my defense in a quite rushing timeframe during summer semester.
I specially acknowledge the support of all my colleagues in our research group.
Although most of us are all working towards different research directions, we established vi
a very strong collaboration and stimulating discussion across different fields. Among
them I thank especially Xuejun Huang, Scott Sutton and Emre Cinkilic for their constant
assistance in most of my critical experimentations, the titanium casting and testing. I
want to mention three of our group’s post-docs, Dr. Weihua Sun, Dr. Renhai Shi, and Dr.
Jiashi Miao, for providing me immense help in CALPHAD database-ing and microstructural characterizations. Also, as mentioned by our previous graduating colleague, Scott Sutton, I must thank to Janet Meier, Emre Cinkilic, and Scott Sutton for our routine “coffee caravan” with lots of “academia-bouts-exchanging” discussions that really gave me plenty of brand new (and sometimes weirdo) ideas to explore.
I acknowledge the MSE Department’s support staff for assisting us in accelerating
our research works. Especially, I would like to appraise Pete Gosser for helping me to
achieve a lot of casting mold machining and resolving the mechanical testing issues that
really saved me in a pressing need.
Finally, I think my family deserves the greatest thanks from me. My parents have
always stood with me through my entire college career in both China and the United
States, especially for their support and understandings for the path I chose during my
PhD period.
vii
Vita
Sep 2008……….B.S. China University of Mining and Technology, Xuzhou, China
Sep 2010……….B.S. University of Kentucky, Lexington, Kentucky, USA
Sep 2013……….Graduate Research Associate, Department of Materials Science and
Engineering, The Ohio State University
Publications
[1] Z. Liang, J. Miao, J.C. Williams, A.A. Luo, Phase transformation and
strengthening mechanisms investigation of a low-cost and high-strength Ti-Al-Fe-
based cast titanium alloy, in preparation.
[2] Z. Liang, J. Miao, R. Shi, J.C. Williams, A.A. Luo, CALPHAD modelling and
experimental assessment of Ti-Al-Mn ternary system, Calphad, Under Review.
[3] Z. Liang, J. Miao, T. Brown, A.K. Sachdev, J.C. Williams, A.A. Luo, A low-cost
and high-strength Ti-Al-Fe-based cast titanium alloy for structural applications,
Scr. Mater., 157 (2018) 124-128.
viii
[4] Z. Liang, W. Sun, A.A. Luo, J.C. Williams, A.K. Sachdev, CALPHAD modelling
and experimental validation of multi-component systems for cast titanium alloy
development, Proceedings of the 13th World Conference on Titanium, TMS, 2016
1937-1941.
Fields of Study
Major Field: Materials Science and Engineering
ix
Table of Contents
Abstract ...... ii
Dedication ...... v
Acknowledgments...... vi
Vita ...... viii
List of Tables ...... xv
List of Figures ...... xvi
Chapter 1 . Introduction ...... 1
1.1 Motivations ...... 1
1.1.1 Titanium alloys ...... 1
1.1.2 Overall concept of the dissertation and project ...... 3
1.2 Organization of the thesis ...... 7
1.3 References ...... 9
x
Chapter 2 . Titanium Alloying and Casting Technologies: State-of-art ...... 11
2.1.1 α stabilizers ...... 12
2.1.2 β stabilizers ...... 14
2.1.3 Neutral and tracing elements ...... 16
2.2 Casting of titanium alloys ...... 19
2.2.1 Melting technologies ...... 20
2.2.2 Casting and molding technologies ...... 24
2.3 References ...... 26
Chapter 3 . Thermodynamic Reassessment of Ti-Al-Mn Ternary System ...... 28
3.1 Introduction ...... 28
3.2 CALPHAD approach and design of experiments ...... 31
3.3 Sample preparation and microstructure characterization ...... 34
3.4 Differential scanning calorimetry ...... 35
3.5 Thermodynamic modelling and parameters optimization ...... 36
3.5.1 Unary phases ...... 36
3.5.2 Solution phases ...... 36
3.5.3 Intermetallic phases ...... 37
3.6 Experimental results...... 39
xi
3.6.1 Microstructure characterization and phase identification ...... 39
3.6.2 Differential scanning calorimetry ...... 46
3.7 Discussions ...... 46
3.7.1 Binary systems ...... 46
3.7.2 Ti-Al-Mn ternary system ...... 47
3.8 Conclusion ...... 59
3.9 References ...... 61
Chapter 4 . Development of Lightweight Casting Titanium Alloy, Ti-6Al-5Fe-0.05B-
0.05C ...... 66
4.1 Introduction ...... 66
4.2 CALPHAD approach and alloy determination ...... 68
4.3 Sample preparation and microstructure characterization ...... 78
4.4 Differential scanning calorimetry ...... 80
4.5 Mechanical property test ...... 80
4.6 Results and discussion ...... 81
4.6.1 Microstructure characterization ...... 81
4.6.2 As-cast tensile property ...... 85
4.6.3 (α + β) and β transus ...... 87
xii
4.7 Conclusion ...... 89
4.8 References ...... 89
Chapter 5 . Manufacture of Lab-scale Prototype Casting Automotive Connecting Rod .. 93
5.1 Introduction ...... 93
5.2 Lab-scale manufacture casting framework setup ...... 94
5.2.1 Basic concept of the framework ...... 94
5.2.2 Equipment and mold design...... 95
5.2.3 OSU facility setup ...... 97
5.3 Casting simulation ...... 100
5.4 Prototype casting in OSU facility ...... 101
5.4.1 Raw materials ...... 101
5.4.2 Melting operation ...... 103
5.4.3 Casting operation ...... 106
5.5 Casting results and post-assessments ...... 106
5.6 Manufacture cost analysis ...... 109
5.6.1 Raw Materials ...... 109
5.6.2 Energy and Labor ...... 111
5.6.3 Equipment ...... 111
xiii
5.6.4 Final Cost Analysis Comparison ...... 111
5.7 References ...... 116
Chapter 6 . Conclusions and Future Perspectives ...... 117
6.1 Conclusions ...... 117
6.2 Ongoing investigations and future perspectives ...... 119
6.2.1 Further investigation of thermodynamic and kinetics of Ti-Al-Fe system .... 119
6.2.2 Further development of the new alloy, T65-0.05BC ...... 122
6.2.3 Involvement of Mn in alloy design ...... 125
6.2.4 Mold coating investigation ...... 125
6.3 References ...... 128
Chapter 7 Supplementary Results ...... 129
Bibliography ...... 130
xiv
List of Tables
Table 3.1 The assessed thermodynamic parameters of Ti-Al-Mn ternary system in this
work. Only assessed parameters are included, the rest of parameters are inherited from
Ref. [3.25]...... 38
Table 3.2 Multiple phase transition temperatures from DSC and ThermoCalc based on assessed database. Temperature is in the unit of °C...... 58
Table 3.3 Design of experiments and measured EDX phase compositions (at.%) ...... 58
Table 4.1 Tensile properties of as-cast Ti-6Al-5Fe-0.05B-0.05C in comparison with as-
cast Ti-6Al-4V [4.16], Ti-6Al-4V-ELI [4.16] and Ti-5Al-2.5Fe [4.17, 4.18]...... 86
Table 5.1 Cost analysis of raw materials ...... 110
Table 5.2 Equipment, energy and labor cost analysis for cast T65-0.05BC ...... 113
Table 5.3 Equipment, energy and labor cost analysis for P/M steel ...... 114
Table 5.4 Cost analysis summary ...... 115
xv
List of Figures
Figure 1.1 Project/Thesis flowchart ...... 6
Figure 2.1 Calculated Ti-Al binary phase diagram. Database: PanTi_2017 ...... 13
Figure 2.2 Calculated (brown) Ti-V, (red) Ti-Fe, and (blue) Ti-Mn phase diagrams.
Database: PanTi_2017 ...... 15
Figure 2.3 Calculated (top to bottom) Ti-O, Ti-N, Ti-B, and Ti-C binary phase diagrams.
Database: PanTi_2017 ...... 17
Figure 2.4 (top) Schematic of a 50 kg VAR-casting setup with (1-5) VAR system, (6-10)
crucible and casting system, and (bottom) schematic of a 1000 kg semi-continuous VAR-
casting system with (1-5) VAR system, (6-11) crucible, casting, and mold controlling
system [2.9] ...... 21
Figure 3.1 Calculated 1000°C isothermal section of Ti-Al-Mn system based on starting
thermodynamic description [3.25] with design of experiments, including phase equilibria
and DSC experiments...... 33
Figure 3.2 Backscatter secondary electron (BSE) SEM image showing the microstructure of Sample 4...... 40
xvi
Figure 3.3 Backscatter secondary electron (BSE) SEM image showing the microstructure of Sample 6...... 41
Figure 3.4 TEM characterization of Sample 9: a) low magnification backscatter
secondary electron (BSE) SEM image, b) high magnification BSE-SEM with three
phases labeled, c) BF-STEM image with inserted selected area diffraction pattern from
Laves C14 phase (along [0001] zone axis) and L12 phase (along [001] zone axis), and d)
BF-STEM image with inserted diffraction patterns from L12 phase (along [001] zone axis
) and L10 phase (along [100] zone axis)...... 42
Figure 3.5 TEM characterization of Sample 11: a) BSE image, b) BF-STEM image of a
lift-out TEM specimen, c) selected area diffraction patterns of β-Ti (along [001] zone
axis), and d) selected area diffraction pattern of Laves C14 phase (along [[21 ̅1 ̅0] zone
axis)...... 43
Figure 3.6 DSC signal and derivative curves of Sample 12 through heating. The dashed lines are applied to determine the onset temperatures...... 45
Figure 3.7 Comparison between the calculated Ti-Al-Mn 1000°C isothermal sections
from (top) base thermodynamic description [3.25] and (bottom) this work with phase
equilibria data from this work, [3.20] and [3.21]...... 49
Figure 3.8 Calculated Ti-Al-Mn 1200°C isothermal section with phase equilibria data
from [3.20]...... 51
xvii
Figure 3.9 Calculated Ti-Al-Mn 1300°C isothermal section with phase equilibria data
from [3.20]...... 52
Figure 3.10 Calculated Ti-Al-Mn 1000°C isothermal section with XRD data from [3.22].
...... 54
Figure 3.11 Calculated Ti-Al-Mn 1000°C isothermal section with XRD data from [3.23].
...... 55
Figure 3.12 Calculated Ti-Al-Mn 1150°C isothermal section with XRD data from [3.24].
...... 56
Figure 3.13 Calculated Ti = 25 at.% isopleth with experimental data from [3.24]...... 57
Figure 4.1 Calculated Ti-(5, 6, 7)Al-xFe isopleths. Database: PanTi_2017 ...... 71
Figure 4.2 Calculated Ti-xAl-(4, 5, 6)Fe isopleths. Database: PanTi_2017 ...... 72
Figure 4.3 Calculated isopleth of Ti-6Al-5Fe-xB-xC. Database: PanTi_2017...... 75
Figure 4.4 Calculated Ti-6Al-5Fe-0.05B-0.05C-xO isopleth. Database: PanTi_2017. .... 76
Figure 4.5 Calculated Ti-6Al-5Fe-0.05B-0.05C equilibrium pathway. Database:
PanTi_2017...... 77
Figure 4.6 Machined ASTM E8 tensile specimen ...... 79
xviii
Figure 4.7 SEM images of as-cast Ti-6Al-5Fe-0.05B-0.05C under different
magnifications ...... 83
Figure 4.8 STEM characterization of the microstructure of Ti-6Al-5Fe-0.05B-0.05C: (a)
Bright field STEM image with a inserted selected area diffraction pattern showing the
Burgers orientation relationship between α and β phase, (b) HAADF-STEM image, (c)
STEM EDS maps, and (d) STEM EDS line...... 84
Figure 4.9 ASTM E8 tensile result of as-cast Ti-6Al-5Fe-0.05B-0.05C ...... 86
Figure 4.10 DSC signal curve of T65-0.05BC through heating. The dashed lines are applied to determine the onset and offset temperatures...... 88
Figure 5.1 (left column) Pre-casting mold with ZrO2 coating of (top to bottom) top half
and bottom half, and (right column) post-casting mold pictures of (top to bottom) top
half, bottom half, and core ...... 98
Figure 5.2 (top) OSU facility ISM-Casting furnace and (bottom) layout in the vacuum
chamber ...... 99
Figure 5.3 EKK simulation of (top row) final temperature/solidification and (bottom
row) filling time ...... 102
Figure 5.4 Schematics of charge material stacking in DI-water-cooled copper crucible
...... 105
xix
Figure 5.5 Prototype casting T65-0.05BC connecting rod ...... 108
Figure 5.6 The effect of P/M stainless steel powder price on final cost estimation ...... 115
Figure 6.1 Experimental 1000°C isothermal section of Ti-Al-Fe system [6.1]...... 121
Figure 6.2 SEM images of heat treated Ti-6Al-xFe: 1100°C/1hr + 800°C/1hr ...... 123
Figure 6.3 TTT experiments of Ti-6Al-5Fe-0.05B-0.05C: β-homogenization at
1100°C/1hr + (row 1) 750°C, (row 2) 800°C, (row 3) 850°C. The aging time is tagged.
...... 124
Figure 6.4 SEM images of Ti-6Al-xFe-yMn: β-homogenization at 1100°C/1hr +
800°C/1hr ...... 126
Figure 6.5 (a) The molten titanium – ceramic couple experiment setup in plasma arc melter, and SEM images of the ZrO2-titanium interfaces of (b) Ti-6Al-2Fe-2Mn and (c)
Ti-6Al-1Fe-3Mn ...... 127
Figure 7.1 SEM images of as-cast Ti-6Al-5Fe-xB-xC: (top) low magnification, and
(bottom) high magnification ...... 129
xx
Chapter 1 . Introduction
1.1 Motivations
1.1.1 Titanium alloys
The element titanium was firstly analyzed and named by a German chemist,
Martin Heinrich Klaproth, in 1795. However, the interest in titanium rose up much later around 1940’s to 1950’s. The investigation on titanium started from producing titanium sponge and then extended to titanium alloy design. The most important advance was the development of Ti-6Al-4V in the United States in 1954, one of the most successful and wide-applied α-β titanium alloys [1.1]. Since Ti-6Al-4V reaches a good combination between properties and manufacture capability, it still dominates the world titanium application field till today.
Due to its weight-savings, fatigue strength, and intermediate temperature performances [1.1-1.4], titanium and titanium alloys have attracted more attention as potential structural materials which require more ideal property-weight ratios. However, the application of titanium alloys is still limited compared with aluminum, magnesium alloys and steel because of its high cost in manufacture, especially for cost-sensitive civil applications. The cost of titanium alloys is generally composed of two aspects: the raw material and processing. The raw material aspect includes the sponge titanium production
1
and the alloying elements. The sponge titanium production is not related to the scope of
this work and therefore not illustrated. As for the alloying elements, the cost can be
further divided based on α, β stabilizers and neutral elements. For α stabilizers, Al is the
most common and economically reasonable, which does not induce the material cost issue. However, several other interstitial α stabilizer, such as oxygen, needs to be controlled to obtain the desired balance between strength and ductility levels, and thus introduces cost in purity control during manufacturing processes. For β stabilizer, the β
isomorphous category is the most widely used but also expensive, such as V, Mo and Ta.
On the other hand, the β eutectoid category is cheaper but introduces problem of segregation and intermetallic phases, which are detrimental to physical properties. For
neutral elements, they are generally expensive and mostly for morphological control,
which are not considered in this specific project but will be investigated in the future if
necessary.
The other important cost factor for titanium alloy is the manufacture process. For
intricate-shape component, machining is the conventional choice for final shaping.
However, titanium and titanium alloys are generally cost-inefficient in machining due to
their low thermal conductivity. The heat generated from machining process is hard to be
evacuated out of the tools, therefore greatly reduces the tool life and increases the
machining cost. Hence, instead of metal-removal shaping, the net-shape and near-net-
shape processes are more preferable for titanium alloys since they maximally avoid the
machining issues and minimize the metal waste during component realization.
2
This work is focused on addressing the cost issue in titanium alloy manufacture by applying appropriate alloy design and manufacture optimization strategies with a combination approach of simulation and experiments, which is explained in the following sub-section.
1.1.2 Overall concept of the dissertation and project
This work was initiated by the corresponding Department of Energy project targeting at reducing the cost of titanium alloy manufacture for automotive applications.
This dissertation covers multiple topics investigated within the scope of the DOE project, including Calculation of Phase Diagram (CALPHAD) for titanium alloy systems,
CALPHAD-assisted titanium alloy design and development for certain application, and design of correlated manufacture process with Integrated Computational Materials
Engineering (ICME) tools. In this 4-year project, we started from exploring new low-cost titanium alloy systems with the guidance of ICME-CALPHAD framework, and determined a promising alloy system, Ti-Al-Fe-Mn-B-C, that can allow further development (lab-scale and industry-scale) in the future. During the new alloy system exploration, considerable amount of scientific work was completed, including
CALPHAD modeling, microstructure morphology, mechanical properties, and processing of the new alloy systems. The relevant ternary titanium alloy systems, Ti-Al-Mn and Ti-
Al-Fe, were examined and reassessed both computationally and experimentally, leading to more comprehensive understanding of these alloy systems. 3
Based on this above study, a new low-cost casting titanium alloy was developed,
Ti-6Al-5Fe-0.05B-0.05C (designated as T65-0.05BC), which has competitive as-cast
strength level with current commercial titanium alloys but relatively low ductility.
However, in this project, since the target application, automotive engine connecting rod,
is a fatigue-limited component and does not require high ductility, this alloy shall suffice as the deliverable alloy in this project. It should be noted that in order to meet the field- test standard, this new alloy still needs to be thoroughly investigated in heat treatment,
which is a routine procedure for casting titanium alloys.
After designing and confirming the potent of the new alloy, we designed and built
a lab-scale manufacturing framework in OSU facility to conceptually prove that this alloy
can be manufactured with permanent mold casting, which is low-cost and potentially a
mass production process. We designed and produced a lab-scale prototype connecting rod
demonstrating the capability of permanent mold casting of this new alloy with ICME tools. Based on the results from lab-scale manufacturing, several recommendations are made for commercialization/industrialization of this process.
Also, in order to address that this alloy and the related manufacturing process we proposed are cost-competitive with the current alloy and manufacturing techniques, a cost analysis was conducted in this investigation. The results showed that compared with current conventional powder metallurgy steel and the Ti-6Al-4V, this new casting titanium alloy has the potent to replace them in actual manufacturing environment.
4
In addition to the major achievements listed above, there are a few issues that
need to be investigated about this new alloy and the manufacturing process. For example, the heat treatment of this new alloy needs to be further investigated to control the stable microstructure to enhance reliability of final components. The permanent mold casting process requires more precise design to reach maximum efficiency during operation.
These issues and challenges will be listed in Chapter 6. The flowchart of this project/thesis is shown in Figure 1.1.
5
Figure 1.1 Project/Thesis flowchart
6
1.2 Organization of the thesis
As briefly described above, the objective of this work is to develop new cost-
effective casting titanium alloys with the approach combining limited experiments and calculations, and to design the related manufacture process for industrial applications.
Several chapters are permitted to include/use of my own published/under-review journal articles by policies of Elsevier [1.5]. Their corresponding references are listed below accordingly. At first, a general review about titanium alloys is summarized in Chapter 2, focusing on alloying element effects, and the state-of-art manufacturing technologies as the background for the alloy and manufacture process design. Necessary details and literature reviews for different topics are expanded respectively in different chapters.
In Chapter 3, Ti-Al-Mn alloy system, one of the sub-systems studied in this project, is taken as an example for assessment of thermodynamic systems with
CALPHAD tool. This chapter involves the design of experiment for CALPHAD assessment, experimental evaluation, and thermodynamic database examination. The other sub-system, Ti-Al-Fe, is still under investigation and the current progress is illustrated. Its remaining issues are explained in Chapter 6. This chapter is a re-use of an under-review journal article that is expected to get published this year [1.6]
In Chapter 4, we propose a new low-cost casting titanium alloy, Ti-6Al-5Fe-
0.05B-0.05C. This chapter includes the design route of this new alloy starting with the determination of alloying elements, getting referential alloy composition by CALPHAD,
examining the new alloy for microstructure features and various properties with different 7
characterization tools. We also attempted to establish the structure-property relationship for this new titanium alloy. This chapter is a re-use of our published journal article [1.7].
In Chapter 5, a lab-scale permanent mold titanium casting framework was setup in OSU facility and casted with the new casting titanium alloy designed in Chapter 4, Ti-
6Al-5Fe-0.05B-0.05C. This chapter includes the concept of this casting framework, the installation of the equipment/accessories, the design of casting process with ICME tools, the detailed melting-casting operations, and finally presenting the conceptual prototype casting. In addition to the design of the manufacture process, a preliminary cost analysis is also conducted to illustrate the cost potent of this new alloy and manufacture process compared with current application in industry, and therefore proves them capable for further commercialization.
The conclusions and discussions on future works are presented in Chapter 6 for further development of the new alloy and manufacture process based on current work.
The investigations in this dissertation are sponsored by the U.S. Department of
Energy under project DE-FOA-0000988. This report was prepared as an account of work sponsored by an agency of the United States Government. Neither the United States
Government nor any agency thereof, nor any of their employees, makes any warranty, express or implied, or assumes any legal liability or responsibility for the accuracy, completeness, or usefulness of any information, apparatus, product, or process disclosed, or represents that its use would not infringe privately owned rights. Reference herein to
8
any specific commercial product, process, or service by trade name, trademark, manufacturer, or otherwise does not necessarily constitute or imply its endorsement, recommendation, or favoring by the United States Government or any agency thereof.
The views and opinions of authors expressed herein do not necessarily state or reflect those of the United States Government or any agency thereof.
1.3 References
[1.1] G. Lütjering, J.C. Williams, Titanium, second ed., Springer, 2007.
[1.2] D. Banerjee and J.C. Williams, “Perspectives on titanium science and
technology,” Acta Mater., 61 (2013) 844–879.
[1.3] J.T. Whittaker, Ductility and Use of Titanium Alloy and Stainless Steel
Aerospace Fasteners, Thesis, University of South Florida, Tampa FL, 2015.
[1.4] D. Eylon, F.H. Froes, R.W. Gardiner, J. Met., 35 (1983) 35–47; also, in:
Titanium Technology: Present Status and Future Trends, F.H. Froes, D. Eylon,
H.B. Bomberger (Eds.), Titanium Development Association, 1985, pp. 35–47.
[1.5] Elsevier. 2017. Copyright. [ONLINE] Available at:
https://www.elsevier.com/about/policies/copyright. [Accessed 15 August 2018].
[1.6] Z. Liang, J. Miao, R. Shi, J.C. Williams, A.A. Luo, CALPHAD Modeling and
Experimental Assessment of Ti-Al-Mn Ternary System, Calphad, under review. 9
[1.7] Z. Liang, J. Miao, T. Brown, A.K. Sachdev, J.C. Williams, A.A. Luo, A Low-
cost and High-strength Ti-Al-Fe-based Cast Titanium Alloy for Structural
Applications, Scr. Mater., 157 (2018) 124-128.
10
Chapter 2 . Titanium Alloying and Casting Technologies: State-of-art
2.1 Alloying effects on casting titanium alloys
In general, the effects of alloying elements on pure titanium are categorized based
on how they affect the phase stabilities of two essential phases in titanium alloys:
hexagonal close-packed (HCP) α and body-centered cubic (BCC) β phases. The phase
fractions and morphologies of these two phases directly determine the anticipated
properties (especially mechanical properties) of titanium alloys. The phase transformation
temperature between α and β phases, the β transus, is at 882°C in pure titanium. Different
alloying elements can either increase or decrease this signature temperature respectively,
and also introduce a (α + β) two phase region in the middle. The methodologies for
different alloying elements are α stabilizer (increase β transus), β stabilizer (decrease β
transus), and neutral element (no obvious effect on β transus). It should be noted that
within the neutral element type, though they do not significantly affect the β transus,
some elements do have other effects, such as microstructural morphology. These
elements will be specifically discussed in 2.1.3 with other neutral elements since they are
also important factors in alloying design. And since this work focuses on titanium
casting, the casting-related characteristics are also discussed.
11
2.1.1 α stabilizers
The most common α stabilizer is aluminum despite from several tracing elements
such as oxygen, nitrogen, and carbon. Aluminum exists in almost every titanium alloys as
the α stabilizer since it is the only element raising β transus and has large solubility in a large temperature-composition space in titanium as shown in Figure 2.1, the Ti-Al binary
phase diagram. There are other alloying elements that can serve as α stabilizers, but due to their low solubility in titanium, they are not discussed in this work. According to the
binary phase diagram in Figure 2.1, aluminum can raise the β transus by ~24°C/wt.%,
and the maximum solubility of aluminum in titanium is 31.11 wt.% at 1503.15°C.
However, for alloy design purpose, the aluminum composition needs to be controlled
below ~15 wt.% due to the increased stability of ordered β phase as shown in the
neighboring two-phase regions, which is experimentally proved in multiple high
aluminum content titanium alloys [2.1]. The increase stability of Ti3Al (α2) phase should also be concerned. This ordered HCP structure precipitates from disordered α phase, and
can potentially strengthen α phase but severely undermines alloy ductility at room
temperature with large volume fraction [2.2]. The strengthening effect from Ti3Al is
usually applied in γ-TiAl alloys but not common in α-β alloys [2.3].
12
T[C]
Figure 2.1 Calculated Ti-Al binary phase diagram. Database: PanTi_2017
13
2.1.2 β stabilizers
The β stabilizers are further categorized into isomorphous and eutectoid ones depending on whether a eutectoid reaction occurs in the binary phase diagram. In this
sub-section, the involved β stabilizers in this work are discussed: vanadium as the
baseline, iron and manganese as potential candidates.
As can be seen from the calculated binary phase diagrams in Figure 2.2,
compared with vanadium, iron and manganese are more effective in reducing β transus,
and allow β phase to be stable at room temperature. More importantly for casting, these
elements suppress melting temperatures (liquidus). For casting process, the potential
decrease in melting temperature for several hundred degrees can significantly lower the
reactivity issues and difficulty in handling the molten titanium [2.4]. However, the
enlarged liquidus-solidus gap is regarded as the cause of solute segregation during
solidification, which in titanium alloys is known as the “β fleck” [2.5]. The β fleck is a
common problem in the ingot production when the segregation leads to different
compositions of β stabilizers across different regions in macro-scale, and consecutively
leads to different post-heat-treatment behaviors since different β stabilizer compositions
have different β transuses. This segregation problem is possible to be suppressed with
faster cooling rate in the metallic mold casting process to prevent solute from diffusion
(compared with the cooling rate and diffusion distance of large ingot).
14
T /C
Figure 2.2 Calculated (brown) Ti-V, (red) Ti-Fe, and (blue) Ti-Mn phase diagrams.
Database: PanTi_2017
15
2.1.3 Neutral and tracing elements
Most neutral elements, such as Zr, Hf, and Sn, are heavy elements compared with
titanium and cost-ineffective. Since they do not meet the requirements in this
work/project, these elements are not discussed. For trace elements like oxygen and
nitrogen, they are usually inevitable to be included during the processing of titanium
alloys and need to be considered. As for boron and carbon, their involvement originates
from multiple literature investigations.
As shown the binary phase diagrams in Figure 2.3, oxygen and nitrogen are more
soluble in titanium compared with boron and carbon, and both are strong α stabilizers,
and they can significantly modify the strength level of the alloy with trace compositions
[2.6, 2.7]. On the contrary, boron and carbon have limited solubility in titanium (carbon has some extent of solubility and can be partially considered as solid solution strengtheners below certain amounts and intend to form high temperature compounds, titanium boride and carbide. There are a few investigations on boron and carbon effects on titanium alloys about the strengthening and grain refinements, which are reviewed in details in Chapter 4. Among these four elements, oxygen is conventionally considered as the most critical in titanium alloys in order to control the strength and ductility of the final component.
16
T /C
2000 Liquid + Liquid 1750
1500
1250
1000 + T /C
750
500
250
+ Ti2N 0 0 0.25 0.5 0.75 1 1.25 1.5 1.75 2 Ti wt.%(N) Continued
Figure 2.3 Calculated (top to bottom) Ti-O, Ti-N, Ti-B, and Ti-C binary phase diagrams.
Database: PanTi_2017
17
Figure 2.3 continued
1750 Liquid
b + Liquid 1500
1250 + TiB
1000 T /C 750 + TiB
500
250
0 0 0.25 0.5 0.75 1 1.25 1.5 1.75 2 Ti wt.%(B)
2000 Liquid
1750 TiC + Liquid
1500
1250 + TiC
1000 + T /C
750
TiC + 500
250
0 0 0.25 0.5 0.75 1 1.25 1.5 1.75 2 Ti wt.%(C)
18
2.2 Casting of titanium alloys
Because of the high price of titanium itself, the total cost of titanium alloys is largely dependent on the technique in processing and forming. Due to the fact that conventional machining of titanium alloys are hard as explained in Chapter 1, the requirement for titanium casting is growing rapidly. Traditional processing of titanium alloys makes small parts and incorporates into the final component. Casting, instead, makes net shape casting component, which reduces the cost of assembling and metal- removing. For example, with same weight, the cost of each casting part is usually 15-
35% lower than that of wrought part, but without great differences in mechanical properties [2.8]. Usually, titanium casting can be of final product after several heat and surface treatments. Titanium shape casting was firstly applied by U.S. Bureau of Mines with high density graphite molds in 1954, and is growing rapidly right now [2.3].
Between 1978 and 1997, the United States titanium casting shipments had a great increase. Then the decrease was due to golf club production, one of the main users of titanium casting alloy, moved to low-cost countries, and thus another increase in 2003 happened with an increasing need of industrial titanium castings. Titanium casting industry is new compared with other metal casting industries (ferrous, aluminum, and magnesium). With the possibilities of casting techniques improvement and cost reduction, titanium casting is expected to extend its application areas quickly, especially in automotive industries, which requires lighter and stronger components to reduce energy costs. Large-scale titanium casting component has already been widely used in
19
aerospace industry. The challenge addressed in this work is mainly the cost requirements
from massive production in civil areas, specifically in automotive industries in this correlated DOE project. These challenges in titanium casting technology are further discussed in the melting and casting/molding methods respectively.
2.2.1 Melting technologies
Currently, the most commonly used process for making titanium ingot is the vacuum arc remelting technique (VAR). This technique uses the cold-compacted target alloy as the electrode, triggers the electric arc onto the hanged electrode, and gradually melts the electrode to form the ingot at the bottom of the crucible. The ingot is usually reverted and remelted for several times to ensure homogeneity. Today, mature VAR processes can reach a relatively large ingot size and weight (100 cm diameter, 10000-
15000 kg). However, when applying this melting technique with casting process, the setup is complex and not spatially preferable for lab-scale, but should be energy-efficient
and applicable for industry-scale. Two typical large-capacity VAR-casting setup
schematics are shown in Figure 2.4.
20
Continued
Figure 2.4 (top) Schematic of a 50 kg VAR-casting setup with (1-5) VAR system, (6-10) crucible and casting system, and (bottom) schematic of a 1000 kg semi-continuous VAR- casting system with (1-5) VAR system, (6-11) crucible, casting, and mold controlling system [2.9]
21
Figure 2.4 Continued
22
The other new melting technique is the cold hearth melting (CHM). The heating source of CHM processes can be plasma arc, electron beam, or induction field. The
former two approaches are known as plasma cold hearth melting (PCHM) and electron
beam cold hearth melting (EBCHM), and the latter one is known as induction skull
melting (ISM). Different from VAR, during the melting process, the contact location
between the molten metal and cold hearth remains a thin solid layer called “skull”,
preventing any contamination or reaction from the hearth. There are a number of
advantages of CHM compared with VAR, including the purity of the molten metal,
charge material preparation, etc. The most important advantage in this work, specifically for casting process, is the easiness in setup of transferring the molten metal into casting mold. As shown in Figure 2.4, conventional VAR-casting setup needs to transfer melted electrodes into a ladle and then into the casting mold, while the CHM process applied in this work, ISM, can directly transfer the charge materials into the mold. This setup is illustrated in details in Chapter 5. The only disadvantage of CHM process on casting is the limited superheat. Due to the continuous fast heat loss from the cold hearth, the superheat that most cold hearth processes can achieve is 50-100 degrees. It significantly decreases the viscosity of the molten metal and thus limits the intricate cavity filling performance in casting. For EBCHM, similar to VAR, the complexity in casting setup is another limitation for specifically for this work/project. On the other hand, ISM-casting setup is more spatially reasonable for lab-scale experiments, which is applied in this work/project and described in Chapter 5.
23
2.2.2 Casting and molding technologies
Besides the high requirements for melting technique, the other big challenge for
titanium casting technology is the high affinity of titanium with impurities such as
oxygen, nitrogen and hydrogen, and the high melting point of titanium, which requires
that titanium casting must be conducted under vacuum or inert-gas-protected condition
and also within special mold. Titanium casting is developing because of three key factors:
improvement of net shape casting technology, improvement of fatigue properties, and
preservation in casting-mold reaction [1.1]. Currently, there are generally two techniques
for titanium casting: conventional casting with rammed graphite mold, and investment
casting.
Conventional gravity casting is a common technique in the family of casting, but
applying rammed graphite mold is specific for titanium casting (or other highly reactive
metals). The graphite mold is usually made from compacted graphite powder and
organic/inorganic binders. The purpose of rammed graphite mold, which is different from
sand mold for ferrous and bronze casting, is to maximally reduce the reaction between molten titanium casting and the mold. Thus, with this decrease in interfacial reactions, rammed graphite titanium casting can produce complex net shape components with good surface conditions (after regular tumble cleaning and chemical milling) [1.1].
Investment casting is another common technique in casting. Compared with
conventional gravity casting, the first difference is the method of making mold. It firstly makes a wax mold cavity pattern, and then applies a nonreactive coating on it in order to 24
prevent reaction between ceramic mold and molten castings (not necessary). After the
coating is finished, with applying mold materials (ceramics, etc.) and removal of inner
wax by heating, the mold is completed. The advantage of investment casting, comparing
with conventional casting, is that it can produce more complex shape while avoiding
over-extra machining. Since the cavity, i.e. shape of casting, is made by wax (or similar
materials), the pattern can be easily shaped into very complicated forms without heavy
machining. But its disadvantage is also originated from its complicated procedure:
relatively high labor cost. So it is vastly applied in aerospace fields for producing large,
complex, and non-massive production of structural components.
In order to combine the advantages of investment casting and conventional
graphite casting – complex casting infrastructure and reusable mold, permanent mold casting is one of those options. Permanent mold casting usually uses metallic mold, which can provide better thermal conductivity, i.e. heat exchange rate to gain faster cooling rate of casting, and thus, better as-cast microstructure. Permanent mold casting are usually applied for casting of relatively lower melting temperature metals, such as aluminum and magnesium, and also relatively less reactive high melting temperature materials, such as steel. And as stated at the very beginning, titanium has a high melting temperature and also highly reactive with regular metal mold materials, which is difficult in applying permanent metal mold casting. Therefore, the key point in applying permanent metal mold casting for titanium alloy is to prevent reactions between mold and
25
molten metal, which can be achieved via two approaches: specialized mold and modification of casting alloys, which are the main contents in this work.
2.3 References
[2.1] K. Das, S. Das, Order-disorder transformation of the body centered cubic phase
in the Ti-Al-X (X=Ta, Nb, or Mo) system, J MATER SCI, 38 (2003) 3995-4002.
[2.2] R. Boyer, E.W. Collings, G. Welsch, Materials Properties Handbook: Titanium
Alloys, ASM International, 1994.
[2.3] J.C. Chesnutt, Titanium Aluminides for Aerospace Applications, in Superalloys,
TMS, 1992.
[2.4] M. Guclu, Titanium and titanium alloy castings, ASM Handbook Volume 15
Casting, ASM International, 2008.
[2.5] A. Mitchell, A. Kawakami, S.L. Cockcroft, Beta fleck and segregation in
titanium alloy ingots, High Temp Mater Proc, 25 (2011) 337-349.
[2.6] I.I. Kornilov, Effect of oxygen on titanium and its alloys, Met Sci Heat Treat+,
15 (1973) 826-829.
[2.7] W.L. Finlay, J.A. Snyder, Effects of three interstitial solutes (nitrogen, oxygen,
and carbon) on the mechanical properties of high-purity, alpha titanium, JOM, 2
26
(1950) 277-286
[2.8] I.J. Polmear, Light Alloys – From Traditional Alloys to Nanocrystals, 4th
Edition, Elsevier, 2006.
[2.9] Vacuum Arc Remelting, ASM Handbook, 15 (2008) 132-138.
27
Chapter 3 . Thermodynamic Reassessment of Ti-Al-Mn Ternary System
3.1 Introduction
Titanium alloys are increasingly important engineering structural materials for weight-savings, fatigue strength, and intermediate temperature performances [1.1, 1.2,
3.1]. The relatively high costs of titanium alloys have limited their applications to critical components, the development of new cost-effective titanium alloys is necessary to extend their high-volume applications. In order to reduce the cost of titanium components, either the alloy composition or the process (or both) can be tailored to address this issue. The properties of engineering alloys are dependent on their phase constitution and microstructure, resulting from the alloy composition and processing conditions. The alloy constituents and their morphologies are dependent on the thermodynamic and kinetic conditions of the alloy system. Therefore, it is essential to acquire the accurate thermodynamic description of the phase equilibria in order to enable the alloy design in titanium alloy systems.
Compared with aluminum, magnesium and steel, titanium is more expensive due to its raw material cost. Naturally, replacing the current expensive alloying elements (V,
Mo, etc.) with inexpensive alternates is an important approach to reduce the alloy cost
28
[1.2]. Thus, it is necessary to design new low-cost titanium alloys with inexpensive
alloying elements.
In this investigation, manganese was chosen as a potential alloying element to reduce the alloy raw material cost. Manganese is a strong β stabilizer in titanium, but also will significantly lower the melting and β transus temperatures as shown in Figure 2.2.
This also reduces the manufacturing difficulties in melting and heat treatment processes in applications such as casting and additive manufacturing (AM). Also, due to its large solubility in β-Ti phase as shown in Figure 2.2, manganese is expected to be a strong solid solution strengthener in titanium alloys. As for the detailed effects of manganese in titanium alloy system, the potential strength and plasticity of Ti-Al-Mn alloys were investigated by Luzhnikov and Metallovedenie [3.1]. A few studies indicate that the
addition of manganese can potentially improve the ductility in γ-TiAl alloys [3.2, 3.3],
and several γ-TiAl alloys with dilute amount of manganese additions show good
manufacturability, high temperature strength, and fatigue properties [3.5, 3.6]. In
addition, according to a number of investigations [3.7-3.14], manganese also showed
good alloying effect in certain dilute α-β (Ti-2Al-1.5Mn, Ti-2.6Al-2.2Mn, Ti-2.5Al-
1.8Mn) and near α (Ti-0.75Al-0.75Mn-0.3Fe) Ti-Al-Mn alloys regarding tensile
properties, weldability, post-welding mechanical properties, and superplastic deformation behavior. A series of α-β and near-α Ti-Al-Mn alloys were reported by Moissev for
balanced mechanical properties during cold and hot deformation [3.15]. All these reports
suggested potential applications for Ti-Al-Mn titanium alloys. Therefore, Ti-Al-Mn
29
phase diagrams should be examined in order to develop potential new Ti-Al-Mn alloys.
Specifically for Ti-rich alloy design, the major phase regions regarding α, β, and Ti3Al are important since β phase normally serves as the matrix phase, α as strengthening phase for β, and Ti3Al phase as strengthening phase for α. Additionally, acquiring accurate
information about γ-TiAl and its neighboring phases can benefit the design for Mn-
containing γ-Ti alloys as well. Overall, the relationship between the phase equilibria and
alloying elements in this system is essential to design the proper alloy compositions for
specific applications.
Although some thermodynamic assessments have been reported for this ternary
system, most experimental results and assessments were focused on γ-TiAl, Ti3Al and
surrounding Laves phase [3.21, 3.25, 3.26, 3.27], including L12 and Laves C14 phases.
The experimental Ti-Al-Mn phase diagrams were reported in a few investigations [3.16-
3.24]. Chen et al. studied the phase equilibria of alloy composition Ti-42Al-10Mn (at.%) at 1000°C and 800°C [3.21]. Butler et al. investigated the phase transformation temperatures of several γ-TiAl based alloys experimentally and compared with computational results [3.17-3.19]. Kainuma et al. [3.20] studied the alloy composition Ti-
45Al-4Mn (at.%) at 1300°C, 1200°C, and 1000°C and constructed a few experimental isothermal sections. Mabuchi and Nakayama et al. investigated the ternary L12
compound, Ti25Mn9Al66, and neighboring phase regions experimentally [3.22-3.24].
Chakrabarti reported an experimental Mn-rich isothermal section at 1000°C [3.26], which
was later experimentally confirmed by Yan et al. [3.27]. It should be noted that some
30
experimental results may not be sufficiently accurate for alloy design and thermodynamic
assessment. For example, the equilibrated alloy experiments of Mabuchi and Nakayama
et al. [3.22-3.24] were conducted using a short time span that the specimens might not
reach full equilibrium conditions. Despite the most recent work by Chen et al. [3.25], the
assessments of Ti-rich regions (Ti > 80 wt.%) were scarce in either experiments or
modeling, which affect phase regions containing α, β, and Ti3Al phases. As illustrated
above, the formation of these phases are critical to the final alloy properties. Therefore, in
this paper, the equilibrated alloy method was used to obtain phase equilibria information
in Ti-Al-Mn ternary system, and its thermodynamic description was assessed, especially
for Ti-rich region based on the experimental results. These outcomes will provide an
important basis of assessing Ti-Al-Mn-based multi-component systems and assisting design of new cost-effective titanium alloys in the future. The content of this chapter has been under review in Calphad and expected to be published in 2018 [1.6].
3.2 CALPHAD approach and design of experiments
In recent years, CALPHAD proves to be a useful tool for the new alloy design and validation in the ICME framework [3.33]. As mentioned above, accurate thermodynamic databases are the essential input for CALPHAD simulation to predict the equilibrium phase stability, driving force for phase transformation etc. However, unlike existed well-established databases for steel, aluminum and magnesium, the current titanium-based commercial databases are lack of validations due to the scarce 31
experimental data. Therefore, the Ti-Al-Mn experimental data generated in this work is
highly beneficial to improve the current titanium databases.
In this work, the Ti-Al-Mn 1000°C isothermal section was selected to cover most
important phase regions in Ti-rich corner. The design of experiments was based on the
thermodynamic description from Chen et al. [3.25]. Multiple compositions were prepared
in the isotherm as shown in Figure 3.1 and listed in Table 3.3 to validate/reassess the
calculated phase equilibria. Since the focus of this investigation is the Ti-rich region, the
design of experiments is determined in the 2-phase and 3-phase regions neighbouring of the target single phase regions
32
Figure 3.1 Calculated 1000°C isothermal section of Ti-Al-Mn system based on starting thermodynamic description [3.25] with design of experiments, including phase equilibria and DSC experiments.
33
3.3 Sample preparation and microstructure characterization
The button specimens were prepared with plasma arc melter (Arc melter, SA-200,
Materials Research Furnaces, Inc.) from high purity elements: titanium (99.99%, Kamis
Inc.), aluminum (99.99%, SCI Engineering), and manganese (99.9%, Alfa Aesar). Each specimen was re-melted four times to ensure homogeneity, and then encapsulated in the quartz tubes back filled with ¼ atmospheric pressure of Ar for equilibrated annealing treatment. Each encapsulation included a small Ta-3W (wt.%) block to absorb remaining trace oxygen. The specimens were annealed in tube furnace for 720 hours at 1000°C to reach phase equilibria, and then water-quenched to room temperature. Samples for microstructure characterization were ground and polished following conventional metallographic specimen preparation procedures. General microstructure observation was conducted in a FEI Apreo Scanning Electron Microscopy (SEM) equipped with an
Energy-Dispersive X-ray Spectroscopy (EDAX) system. Cross-sectional Transmission
Electron Microscopy (TEM) specimens for phase identification were prepared in a FEI
Helios focused ion beam (FIB) microscope using lift-out method. All FIB TEM specimens were further cleaned using Fischione 1040 nanomill to reduce ion beam damage associated with FIB sample preparation. TEM characterization was carried out on a FEI Tecnai TF20 TEM/STEM microscope operating at an accelerating voltage of
200keV.
34
3.4 Differential scanning calorimetry
The differential scanning calorimetry (DSC) was conducted with the NETZSCH
DSC 404F1 Pegasus® in this investigation at University of Science & Technology
Beijing in order to evaluate the 3-phase region, (β-TiMn + β + Laves C14), which is hard
to pinpoint via equilibrated alloy method. The DSC specimen was prepared with the
equilibrated alloy method above. The specimen, Ti-2.24Al-48.94Mn (wt.%), was placed
in an alumina crucible with an empty alumina crucible as reference. The specimen
chamber is evacuated and backfilled with argon gas. The empty alumina crucible was
tested before the sample to obtain the baseline, which was deducted for the results of the sample. The DSC instrument, namely the thermocouple, has been calibrated regularly.
The calibration process includes both the temperature and sensitivity calibrations, which are specific to the crucible, the heating rate, and the protective gas. Five pure metals, i.e.
Au, Al, Zn, Bi, and In, were applied for the calibrations. The temperature calibration
eliminates the deviation between the temperature measured by the thermocouple and the
actual temperature of the sample. Sensitivity correction ensures the accurate conversion
between thermocouple signal and heat flow power. The chamber temperature was
increased to 1200°C at the rate of 20°C min.
35
3.5 Thermodynamic modelling and parameters optimization
3.5.1 Unary phases
The Gibbs free energy function for the element Ti, Al, Mn in all phases is described by Eq. (1) from SGTE compilation of Dinsdale [3.30]:
⁄ ⋯ (1)
where ϕ represents the target phase, and T represents the absolute temperature.
3.5.2 Solution phases
In the Ti-rich region of Ti-Al-Mn ternary system, the solid solution phases liquid,
BCC, and HCP phases are described by Eq. (2):
ln , ,
, , , ,
(2)
The first two terms represents ideal solution, and the third and the fourth terms
represents the excess Gibbs free energy. xi is the mole fraction of component i. is taken from Eq. (1). R is the gas constant. L’s are binary and ternary interaction parameters in the third and the fourth terms. Vk is defined
as: 1 ∑ , , ⁄ [3.22].
36
3.5.3 Intermetallic phases
Multiple intermetallic phases in the Ti-Al-Mn ternary system, including Lave
C14, α2 (Ti3Al), γ (TiAl), Al8Mn5, and L12, are described by Eq. (3):
(3)
The thermodynamic descriptions from Chen et al. [3.25] were taken as the pre- assessment model and re-assessed for Ti-rich region.
In this work, all thermodynamic parameters are optimized manually in
ThermoCalc. The optimization priority of experimental data is as follows: three phase tielines, two phase tielines, and XRD data points. All pre- and post-optimized parameters are listed in Table 3.1.
37
Table 3.1 The assessed thermodynamic parameters of Ti-Al-Mn ternary system in this
work. Only assessed parameters are included, the rest of parameters are inherited from
Ref. [3.25].
Phase Parameter Liquid , , 20000 18.33 α , , 28427.33 26.334 , , 46657.5 50 β , , 81154.51 35.4668 , , 29460 26.8 Laves C14 _ _ : 94500 18 2 : 62485 17.5 _ _ 2 , : 6000 , : 135859 66.66 , : 34190 8 _ _ Ti3Al : 350000 3 _ _ : 15000 3 , : , : 60460 21 , : 27000 3.5 L12_Ti25Mn9Al66 , : : 33052 13 , : : 8000 5 : : , 16217 19.67 : : , 2000 : : , 1316.47 5
38
3.6 Experimental results
3.6.1 Microstructure characterization and phase identification
Selected annealed microstructures and diffraction patterns are shown for the assessed multi-phase regions and related phase identifications. Figure 3.2 is BSE-SEM image of Ti-17.75Al-5.62Mn alloy located in phase region (β + Ti3Al). Figure 3.3 shows the microstructure of Ti-27.35Al-7.17Mn alloy located in phase region (Ti3Al + L10 + β).
The compound phases in these two specimens were identified by the composition ratios and did not need further diffraction patterning.
Figure 3.4 a) and b) are BSE-SEM images of Ti-55Al-14Mn specimen. Based on the image contrast, it can be seen that there are three different phases in this alloy. Two site-specific FIB specimens were prepared for phase identifications. Figure 3.4 c) and d) are bright field scanning transmission electron microscopy (BF-STEM) images of these two specimens. Selected area diffraction (SAD) patterns from all phases are shown in
Figure 3.4. The diffraction results show three phases: Laves C14, L10 and L12 phases,
indicating that this alloy is located in a three phase region and consistent with the
CALPHAD simulation. The needle-morphology of the L12 phase is also consistent with
previous SEM investigations by Mabuchi and Nakayama et al. [3.22-3.24]. The
diffraction pattern represents a cubic crystal structure, and the measured lattice parameter
is 0.3727 nm. As for L10 FCC tetragonal phase, the measured lattice parameters a and c
are respectively 0.3953 nm and 0.3980 nm, indicating that the solid solutioning of Mn in
L10 does not significantly increase the tetragonality.
39
Figure 3.2 Backscatter secondary electron (BSE) SEM image showing the microstructure of Sample 4.
40
Figure 3.3 Backscatter secondary electron (BSE) SEM image showing the microstructure of Sample 6.
41
Figure 3.4 TEM characterization of Sample 9: a) low magnification backscatter secondary electron (BSE) SEM image, b) high magnification BSE-SEM with three phases labeled, c) BF-STEM image with inserted selected area diffraction pattern from
Laves C14 phase (along [0001] zone axis) and L12 phase (along [001] zone axis), and d)
BF-STEM image with inserted diffraction patterns from L12 phase (along [001] zone axis
) and L10 phase (along [100] zone axis).
42
Figure 3.5 TEM characterization of Sample 11: a) BSE image, b) BF-STEM image of a lift-out TEM specimen, c) selected area diffraction patterns of β-Ti (along [001] zone axis), and d) selected area diffraction pattern of Laves C14 phase (along [[21 ̅1 ̅0] zone axis).
43
Figure 3.5 a) is a BSE-SEM image of Ti-20Al-32Mn alloy. It can be seen that the
microstructure of this alloy mainly consists of two phases. Figure 6 b) is a BF-STEM
image of a FIB lift-out specimen prepared from this alloy. The electron diffraction
results in Figure 6 c) and d) show that these two phases are ordered-β and Laves C14
phase respectively, in which the ordering of β-Ti is consistent with high Al contained
ordered-β phase in Ti-Al binary system. The measured lattice parameters a and c are
respectively 0.7813 nm and 0.4876 nm, indicating that Al in solid solution of Laves C14 phase compresses the HCP lattice compared with the original one (TiMn2 Laves C14 in
Ti-Mn binary system) [3.34].
There was no unidentifiably phase in experiments. Multiple EDXS point analyses
(>5) were performed for each phase. The phase compositional analysis results were summarized in Table 3.3 and used to plot experimental isothermal sections in section 3.7.
44
Figure 3.6 DSC signal and derivative curves of Sample 12 through heating. The dashed lines are applied to determine the onset temperatures.
45
3.6.2 Differential scanning calorimetry
The DSC result for Ti-2.24Al-48.94Mn (at.%) alloy is shown in Figure 3.6. The
two curves represent DSC heat flow signal and the correlated derivative of temperature respectively. As can be seen, an endothermic peak and an exothermic peak appear during
heating, indicating two phase transition temperatures. The first transformation
temperature is considered as the onset point in the DSC signal curve defined by the
intersection of the tangential lines between inflection point and plateau, which is 857°C.
But since the second transformation is lack of the plateau to determine the onset point, the inflection point to estimate the boundary between two peaks, and therefore can be regarded closer to the actual phase transition temperature, which is 968°C. And at the end of the signal curve, a sharp endothermic drop is observed indicating that the solidus has been reached and the specimen starts to melt. The starting temperature of the endothermic
drop is considered as the solidus of the specimen, which is 1188°C. The DSC result will
be used to assess and compare with the thermodynamic description in section 3.7.
3.7 Discussions
3.7.1 Binary systems
As mentioned by Chen et al [3.25], three binary systems in the Ti-Al-Mn system
have already been fully assessed by Witusiewicz, Du, and Saunders [3.28-3.30] with the
most recent experimental data. Therefore, in this work, all binary parameters in the
46
thermodynamic model are inherited and not reassessed. The assessment in this work includes only the modification of current ternary parameters.
3.7.2 Ti-Al-Mn ternary system
The phase equilibria in the Ti-Al-Mn ternary system at 1000°C, 1150°C, 1200°C, and 1300°C have been examined based on equilibrated alloy experiments in this
investigation and literature [3.18, 3.20-3.24]. The experimental results from literature
included phase equilibria, DTA, and XRD. It should be noted that not all literature data is used in the assessment due to the reasons discussed in section 3.1. The literature experimental results with less confidence are only applied as indirect references to check the post-assessed results. Figure 3.7 to Figure 3.12 show the recalculated isothermal sections with the experimental data at 1000°C, 1150°C, 1200°C, and 1300°C respectively
(Figure 3.7 includes comparison between [3.25] and this work). With the experimental data from this investigation, it can be found that compared with the starting thermodynamic description [3.25], a better agreement has been reached for multiple two- phase and three-phase regions between α, β, L10, Ti3Al, Laves C14, and L12. The
assessment in this investigation specially focuses on the three phase regions (β + L10 +
Laves C14), (β + L10 + Ti3Al), and (Laves C14 + L10 + L12). It should be noted before
the discussion that the XRD experimental data from references [3.22-3.24] are considered
less accurate due to their experimental methods described in the literature. Their
homogenization and heat treatment times are too short to reach phase equilibria. 47
Therefore, in the assessments, the experimental data from references [3.18, 3.20, 3.21] are considered of preference, while the ones from references [3.22-3.24] are only applied to check the relative accuracy of the assessments.
48
[3.21] and this work
[3.20]
Continued
Figure 3.7 Comparison between the calculated Ti-Al-Mn 1000°C isothermal sections from (a) base thermodynamic description [3.25] and (b) this work with phase equilibria data from this work, [3.20] and [3.21].
49
Figure 3.7 continued
[3.21] and this work
[3.20]
50
[3.20]
Figure 3.8 Calculated Ti-Al-Mn 1200°C isothermal section with phase equilibria data from [3.20].
51
[3.20]
Figure 3.9 Calculated Ti-Al-Mn 1300°C isothermal section with phase equilibria data from [3.20].
52
Figure 3.7 to Figure 3.9 show the calculated isothermal sections of 1000°C,
1200°C, and 1300°C with the experimental results based on the phase equilibria method.
It can be seen that the three-phase regions, (β + L10 + Laves C14), (β + L10 + Ti3Al), and
(β + Laves C14 + L12), are given with high priority to fit the experimental data. As
shown in Figure 3.7, the experimental result of (β + L10 + Laves C14) in this work
closely resembles the experimental data from reference [3.21], which also proves the
reliability of the phase equilibria method. Especially, as compared in Figure 3.7 between
the thermodynamic description in reference [3.25] and this work, all phases related to Ti- rich region are precisely predicted in this work.
Figure 3.10 to Figure 3.12 show the calculated isothermal sections of 1000°C and
1150°C with the XRD data from literature. It can be seen that, compared with the thermodynamic description in reference [3.25], the assessed one is more consistent for most three-phase data points, which are all associated with Laves C14 and L12 phases.
However, as mentioned above, due to the related experimental procedures described in the literature, the XRD data is used for comparison, not for performing the assessment.
53
[3.22]
Figure 3.10 Calculated Ti-Al-Mn 1000°C isothermal section with XRD data from [3.22].
54
[3.23]
Figure 3.11 Calculated Ti-Al-Mn 1000°C isothermal section with XRD data from [3.23].
55
[3.24]
Figure 3.12 Calculated Ti-Al-Mn 1150°C isothermal section with XRD data from [3.24].
56
[3.24]
Figure 3.13 Calculated Ti = 25 at.% isopleth with experimental data from [3.24].
57
Table 3.2 Multiple phase transition temperatures from DSC and ThermoCalc based on assessed database. Temperature is in the unit of °C.
Phase Transition Phase Transition Phase
region 1 Temp. 1 region 2 Temp. 2 region 3 α-TiMn + β DSC 856.703 β-TiMn + β + 967.917 β + + Laves Laves C14 Laves C14 Calculated C14 867.834 998.640
Transition Phase
Temp. 3 region 4 DSC 1188.000 Liquid+ β Calculated 1200.300
Table 3.3 Design of experiments and measured EDX phase compositions (at.%)
Measured Phase 1 Phase 2 Phase 3 Point specimen Al Mn Al Mn Al Mn Al Mn 1 23.11 2.33 25.50 0.00 20.06 4.41 N/A 2 23.02 2.33 25.42 0.00 19.66 4.41 N/A 3 27.88 4.34 29.04 2.37 26.05 8.43 N/A 4 34.96 4.08 34.09 9.73 34.91 4.09 N/A 5 39.85 17.07 35.37 30.93 34.85 13.67 45.74 5.12 6 40.36 5.19 46.00 3.00 35.20 7.80 36.90 3.60 7 51.50 9.30 43.90 23.30 54.30 4.50 N/A 8 21.30 30.40 18.20 43.80 23.67 19.40 N/A 9 56.10 12.80 43.50 23.40 54.02 6.00 60.06 12.04 10 18.83 37.08 16.11 46.48 22.49 20.71 N/A 11 22.37 31.28 18.82 44.64 24.46 19.83 N/A * This specimen is for DSC experiment and does not 12 2.24 48.94 have phase composition measurements
58
Figure 3.13 shows the assessed 25 at.% Ti isopleth with the experimental data
[3.24]. It can be seen that except for the liquidus temperature, the calculated result is
sufficiently accurate. It should also be noted that according to Mabuchi et al. [3.24], the
liquidus was taken from the peak temperature of DTA signals (no DTA curves were
reported in the original article). In order to assess the liquidus, the endpoint temperatures
should be considered instead of peak temperatures. Therefore, in reference 3.24, the
liquidus temperatures might be underestimated.
Table 3.2 compares the DSC results with calculated phase transition temperatures
for Ti-2.24Al-48.94Mn. The results show good agreement between the experimental data
and calculations from the assessed database, giving credibility to the accuracy of the 3-
phase region (β + L10 + Laves C14) in the assessed database.
3.8 Conclusion
Based on the previous thermodynamic assessments for the Ti-Al-Mn ternary
system, the thermodynamic description is accurately reassessed specifically for the Ti-
rich region depending on the phase equilibria experimental results. These combinatory
experimental and thermodynamic assessments have provided reliable CALPHAD results
for designing Ti-Al-Mn-based new cost-effective titanium alloys. Based on the
experimental-based assessments, the following major achievements have been drawn:
59
(1) The Ti-Al-Mn ternary system is refined based on the latest experimental results,
especially with new experimental data for the Ti-rich corner from this work,
which is valuable for updating the current commercial titanium thermodynamic
databases.
(2) The presence of L12 phase is directly proved with both microstructure
morphology and crystal structure observations, including the precisely measured
lattice parameter a = 0.3723 nm. The TEM work in this study clearly shows its
cubic crystal structure.
(3) The maximum solubility limits of aluminum and manganese in β-Ti phase region
are modified, indicating a lower β-Ti phase stability limit with increased amount
of aluminum and manganese. These alternations are important for maximizing
alloying effects in Ti-rich Ti-Al-Mn alloy design.
(4) The maximum solubility limits of manganese in Ti3Al and (α + Ti3Al) phase
regions are modified, showing a higher Ti3Al phase stability with increased
manganese content. This alternation is important for designing both titanium
alloys with dilute-Mn content and γ-TiAl alloys. While the former could possibly
introduce certain amount of Ti3Al for precipitation strengthening, and the latter
applies Ti3Al as the main strengthening phase.
(5) The updated thermodynamic description of the Ti-Al-Mn ternary system will be
essential for the determination of Ti-Al-Mn-based n-ary (>=4) thermodynamic
systems for multicomponent low cost titanium alloy design.
60
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Fe additions on the microstructure and mechanical properties of a Ti-Al-Mn
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[3.16] R.F. Domagala, W. Rostoker, The system titanium-aluminum-manganese, Trans
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[3.17] C.J. Butler D.G. McCartnery, C.J. Small, F.J. Horrocks, N. Saunders,
Solidification microstructures and calculated phase equilibria in the Ti-Al-Mn
system, Acta Mater, 45 (1997) 2931-2947.
[3.18] C.J. Butler D.G. McCartnery, An experimental study of phase transformations
and a comparison with calculated phase equilibria in Ti-Al-Mn alloys, Acta
Mater, 46 (1998) 1875-1886.
[3.19] C.J. Butler D.G. McCartnery, Phase transformations and phase equilibria in a Ti-
37%Al-20%Mn alloy, Intermetallics, 7 (1997) 663-669.
[3.20] R. Kainuma, Y. Fujita, H. Mitsui, I. Ohnuma, K. Ishida, Phase equilibria α(hcp),
β(bcc) and γ(L10) phase in Ti-Al base ternary alloys, Intermetallics, 8 (2000)
855-867.
[3.21] Z. Chen, I.P. Jones, C.J. Small, Laves phase in Ti-42Al-10Mn alloy, Scr Mater,
35 (1996) 23-27.
[3.22] H. Mabuchi, K.I. Hirukawa, Y. Nakayama, Formation of structural L12
compounds in TiAl3-base alloys containing Mn, Scr Metall, 23 (1989) 1761-
1766. 63
[3.23] Y. Nakayama, H. Mabuchi, Formation of ternary L12 compounds in Al3Ti-base
alloys, Intermetallics, 1 (1993) 41-48.
[3.24] H. Mabuchi, A. Kito, M. Nakamoto, H. Tsuda, Y. Nakayama, Effects of
manganese on the L12 compound formation in Al3Ti-base alloys, Intermetallics,
4 (1996) 193-199.
[3.25] L.Y. Chen, C.H. Li, A.T. Qiu, X.G. Lu, W.Z. Ding, Q.D. Zhong, Calculation of
phase equilibria in Ti-Al-Mn ternary system involving a new ternary
intermetallic compound, Intermetallics, 18 (2010) 2229-2237.
[3.26] D.J. Chakrabarti, Phase stability in ternary systems of transition elements with
aluminum, Metall Trans B, 8B (1977) 121-123.
[3.27] X. Yan, X. Chen, A. Grytsiv, P. Rogl, R. Podloucky, H. Schmidt, G. Giester, X.
Ding, On the ternary Laves phases Ti(Mn1-xAlx)2 with MgZn2-type,
Intermetallics, 16 (2008) 16-26.
[3.28] V.T. Witusiewicz, A.A. Bondar, U. Hecht, S. Rex, T.Y. Velikanova, The Al-B-
Nb-Ti system III. Thermodynamic re-evaluation of the constituent binary system
Al-Ti, J Alloys Compd, 465 (2008) 64-77.
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64
[3.30] I. Ansara, A.T. Dinsdale, M.H. Rand (Eds.), COST507, Thermochemical
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[3.32] D. Fruchart, J.L. Soubeyroux, R. Hempelmann, Neutron diffraction in Ti1.2Mn1.8
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65
Chapter 4 . Development of Lightweight Casting Titanium Alloy, Ti-6Al-5Fe-0.05B- 0.05C
4.1 Introduction
Titanium alloys started to attract interests as potential lightweight structural materials due to the intermediate temperature properties since 1960’s [1.1]. The relatively high weight-specific properties make titanium alloys ideal for weight-sensitive applications such as aeronautics engines [1.2]. However, its application in civil area is still scarce because of the high manufacturing cost. Currently, Ti-6Al-4V (wt.%) alloy dominates about 50% of titanium products including castings [1.3], although it was designed for wrought applications only but not optimal for other manufacturing processes. In this investigation, the net-shape and near-net-shape processes are considered. Compared with wrought process, these techniques have cost advantages in producing complex geometry components due to the exclusion or limited machining step as briefly mentioned in Chapter 1, which costs significantly for titanium alloys.
Therefore, it is necessary to design new low cost casting titanium alloys, which is the target in this work.
As mentioned above, net-shape and near-net-shape processes are preferable for titanium alloys by avoiding metal removal cost [4.1]. As concerned of properties, titanium castings can potentially be competitive in mechanical properties with wrought 66
products at lower costs. For instance, Ti-6Al-4V has approximately the same monotonic loading strength for both casting and wrought-annealed products. But for crack growth
resistance and creep strength, wrought products can, while casting cannot, be further tailored by thermomechanical processing at additional cost. Thus, titanium castings are usually applied in structural applications without the requirements of dynamic but static properties, such as yield strength, ultimate tensile strength and toughness [1.4, 4.2]. But on the other hand, casting can be heat-treated to modify strength and ductility by engineering the microstructures. In this chapter, a new casting titanium alloy was developed for structural applications and addressed two challenges in titanium casting: castability and raw material costs.
In this investigation, Ti-6Al-4V was used as the baseline and comparison since it is currently the most widely applied titanium alloys in industry, including the casting, and well-documented for properties and microstructure characterizations as references. The raw material cost was addressed by changing the β stabilizing element vanadium to iron, the strongest β stabilizer and the cost-effective option. Ti-Fe system can also potentially decrease liquidus temperature and enlarge the liquidus-solidus gap as shown in Figure
2.2, which is favorable for casting processes. However, in the meantime, the introduction of iron will lead to significant solute segregation as well due to the eutectic nature of Ti-
Fe system, known as β fleck in titanium alloy, but might be alleviated by fast cooling rate in permanent mold casting. In this chapter, a combinatory experimental and CALPHAD approach was applied to develop a new Ti-Al-Fe alloy with both experimental and
67
thermodynamic basis in the as-cast condition. The content of this chapter has been published in Scripta Materialia [1.7].
4.2 CALPHAD approach and alloy determination
Multiple isopleths in Ti-Al-Fe system were calculated in order to get the phase equilibrium information with certain aluminum or iron compositions to determine the desired phases and their volume fractions. Since the goal of this investigation/project focuses on the cost reduction, aluminum and iron contents are ideally supposed to be included as much as possible to compensate the titanium cost. However, the involvement of aluminum and iron will increase the phase stabilities of three embrittlement phases, ordered β-Ti, Ti3Al and TiFe, which is the main restricting factor for their maximal
thresholds.
As shown in multiple isopleths with different aluminum contents in Figure 4.1,
the phase stability and relationship among α, β, and TiFe remains the same until
aluminum content reaches 7 wt.%. At 7 wt.% of aluminum, the Ti3Al starts to form at
intermediate temperature (~550°C). Since Ti3Al is not considered in this project due to its
brittle nature, aluminum content is preliminarily controlled at 6 wt.% to minimize its
formation. Multiple experimental investigations from literature also lead to the same
trend of aluminum alloying effect in titanium alloy, including the ordering of β-Ti into
B2 structure and the stability of Ti3Al. The ordering of β-Ti is shown in Figure 2.1 above
68
15 wt.% of aluminum in Ti-Al binary system, which should also be avoided due to its brittleness. This ordering phenomenon is observed at much lower aluminum compositions in Ti-V-Cr β-titanium alloy by Li et al. [4.3 and 4.4]. This order-disorder phase transformation was found at ~620°C with only 4 wt.% of aluminum in Ti-32V-
15Cr-5.6Al. As for the Ti3Al, similar approach was applied on Ti-25V-15Cr alloy, and the Ti3Al formation was observed within 2-4 wt.% by Li et al. [4.5].
As for iron, the same isopleth approach is applied as shown in Figure 4.2. The increase of iron content significantly increase the phase stability of TiFe, reflecting by the disappearance of α single phase region and its replacement of (α + TiFe) two phase region. Especially for the 6 wt.% aluminum content determined above, the disappearance of α single phase region will lead to the direct precipitation pathway of TiFe from primary β matrix and should be avoided. Therefore, the iron content is preliminarily controlled at 5 wt.%. As briefly illustrated in 2.1.2, iron is usually considered difficult to be involved in titanium alloys, especially for ingot metallurgy, due to its strong segregation effects in macro scale. However, in direct casting (from raw charge materials to casting cavity-fill without pre-alloying ingots) for small geometrical width/thickness, the solidification segregation is expected to be suppressed by fast cooling rate. Therefore, iron becomes one of the acceptable alloying elements. The alloying effect of iron in titanium alloy was carried during 1980-2000s. In 1980-90s, several α-β iron-contained titanium alloys were developed for hot working processes, such as Ti-2.5Al-5Fe [4.6],
TIMETAL 62S [4.7] and Ti-5.5Al-1Fe [4.8]. The Ti-Al-Fe system showed some potent
69
as the replacement of Ti-Al-V, but was still challengeable due to the segregation problem
mentioned above, which is common for eutectic β stabilizing elements. Fuji and
Takahashi [4.8] did a relatively comprehensive experimental investigation of the
mechanical properties on multiple Ti-Al-Fe alloys with thermomechanical treatments.
Iron was also investigated as addition for a few biomedical β titanium alloys, including
Ti-Mo-Fe [4.9] and Ti-Nb-Fe [4.10] systems. Phase analyses from both investigations
indicated that iron content promotes the formation of metastable ω phase in β titanium alloys and improvements in bending properties. However, both investigations are limited at experiments without further theoretical discussions for the iron effect. Despite of these experimental investigations above, there is no systematic investigation on Ti-Al-Fe system for the alloying behavior, thermodynamic descriptions, and thus no further development and applications of this alloy system.
70
1700 5Al 1600 Liquid 6Al 7Al 1500 Liquid + 1400
1300
1200
1100 T /C 1000
900 + TiFe 800 + + + TiFe 700
600
+ Ti Al (7Al) + Ti Al + TiFe (7Al) 500 3 3 01234567891011121314151617181920 wt.%(Fe)
Figure 4.1 Calculated Ti-(5, 6, 7)Al-xFe isopleths. Database: PanTi_2017
71
1700 Liquid 1600 Liquid + 1500 4Fe 5Fe 1400 6Fe
1300 + Ti3Al + 1200
1100
T /C + Ti Al 1000 3
900
+ Ti3Al + TiFe 800 +
Ti3Al + TiFe 700 (4Fe, 5Fe) + Ti3Al (4Fe, 5Fe) 600 + TiFe + Ti3Al + TiFe 500 01234567891011121314151617181920 wt.%(Al)
Figure 4.2 Calculated Ti-xAl-(4, 5, 6)Fe isopleths. Database: PanTi_2017
72
Boron and carbon were also involved for grain refinement and α-strengthening respectively. In multiple reported investigations [4.11-4.14], trace levels of boron and carbon were also included in this design to provide grain refinements and the solid solution strengthening effects since casting alloy usually contains large grain size that suppresses alloy properties compared with hot/cold-worked component. The significant grain refinement effect of boron in as-cast titanium alloys was investigated by Roy et al. in the baseline alloy Ti-6Al-4V [4.11], and in multiple aerospace β titanium alloys by
Tamirisadandala et al. [4.12], both showing refinement of primary α and β phases. As for the combined effect of boron and carbon, Banoth et al. [4.13] and Sarkar et al. [4.14] investigated this in multiple metastable β titanium alloys, which suggested that boron and carbon refined primary β grain and the following α platelet sizes correspondingly. Since grain refinement is in general favorable to improve mechanical properties [4.15], boron and carbon are included in trace levels (0.05 wt.% for both based on literatures) in this casting alloy design. The calculated phase equilibrium pathway of trace amount (wt.% for both boron and carbon) was shown in Figure 4.3. As can be seen, 0.05 wt.% of boron and carbon do not modify the phase stability by much in Ti-6Al-5Fe system except for introducing TiC and TiB precipitates. However, it is expected that they will modify the as-cast microstructure morphology. Judging from their formation sequences on the equilibrium pathway, TiB is likely to form with primary β grain and locate on grain boundaries because of the low solubility of boron in β-Ti. While TiC is not likely to form at low carbon content because of the relatively higher solubility of carbon in both α-Ti
73
and β-Ti, carbon might precipitate as TiC under some cases, but will more probably stay in α-Ti or β-Ti as solid solution strengthener. Since the boron and carbon effects are still scarce in experiments, different boron and carbon additions specimens were prepared to investigate as-cast microstructures. The as-cast microstructures show heavy aggregation of large intermetallic phases on the grain/dendritic boundaries with large boron and carbon additions in Figure 7.1. Therefore, in this alloy design, boron and carbon concentrations are remained low. These results are not strictly related to the alloy design in this chapter and sorted in the supplementary results in Chapter 7.
Another concern in titanium alloy design is the oxygen effect. As illustrated in
2.1.3, oxygen is an α stabilizer and can significantly harden α phase. Since oxygen cannot be avoided due to its strong adhesion on titanium, it needs to be evaluated. As shown in
Figure 4.4, the Ti-6Al-5Fe-0.05B-0.05C-xO isopleth is calculated. Oxygen does not show a strong effect on phase transformation temperatures in this system. However, it is expected that oxygen will affect the alloy strength level [1.1]. Therefore, the oxygen level is measured in the experiment to evaluate the oxygen contribution on the strength.
Based on the alloy design above, the equilibrium pathway of the nominal composition is calculated as shown in Figure 4.5. All phases in the equilibria are possible to appear in as-cast status depending on the specific kinetic conditions.
74
1700 Liquid 1600 Liquid + 1500
1400
1300 + TiB 1200
1100 + TiC + TiB T /C 1000
900 + + + TiB + TiC + + TiB 800
700 + TiB + TiC 600 + TiFe + TiC + TiB
500 0 0.02 0.04 0.06 0.08 0.1 0.12 0.14 0.16 0.18 0.2 wt.%(B, C)
Figure 4.3 Calculated isopleth of Ti-6Al-5Fe-xB-xC. Database: PanTi_2017.
75
1700 Liquid 1600 Liquid + 1500
1400
1300 + TiB 1200 T /C 1100
1000 + TiC + TiB 900 + + TiC + TiB 800
700 + TiFe + TiC + TiB 0 0.02 0.04 0.06 0.08 0.1 0.12 0.14 0.16 0.18 0.2 wt.%(O)
Figure 4.4 Calculated Ti-6Al-5Fe-0.05B-0.05C-xO isopleth. Database: PanTi_2017.
76
1
0.9
0.8
0.7
-Ti 0.6 TiC TiFe 0.5 -Ti Liquid 0.4 TiB
0.3
0.2 Volume Fraction of Phases 0.1
0 500 600 700 800 900 1000 1100 1200 1300 1400 1500 1600 1700 T /C
Figure 4.5 Calculated Ti-6Al-5Fe-0.05B-0.05C equilibrium pathway. Database:
PanTi_2017.
77
4.3 Sample preparation and microstructure characterization
The specimens with nominal composition Ti-6Al-5Fe-0.05B-0.05C (designated as
T65-0.05BC, wt.%) were prepared by induction skull melter under high purity Ar
atmosphere (ISM, PVT Inc., an InductoTherm Group company) using CP-2 titanium
(99.9%, ATI, Inc.), CP aluminum (99.9%, Alfa Aesar), CP iron (99.9%, Alfa Aesar),
TiB2 (99.5%, Alfa Aesar), and carbon slug (99.5%, Alfa Aesar). Each specimen was
flipped and re-melted four three times to ensure homogeneity. The oxygen content was
tested by Luvak Inc. with ASTM standard E1409-13. Microstructure characterization and
compositional analysis were conducted using Scanning Electron Microscopy (SEM),
Transmission Electron Microscopy (TEM), and Energy Dispersive Spectroscopy (EDS).
SEM experiments were conducted in an FEI Apreo microscope, operated at 30 kV.
Blanks for TEM foils were sectioned from as-cast specimens using a low speed diamond saw. The foils were mechanically ground to a thickness of about 100 µm. Final thinning and perforation of the TEM foils were completed by electro-polishing in an electrolyte containing 5% perchloric acid, 35% butyl cellosolve and 55% methanol at a temperature of -30°C and an applied voltage of 30 V. TEM characterization was performed on an image-corrected and monochromated Titan3 60-300 microscope equipped with a SuperX
EDS system and operated at 300 KeV.
78
Figure 4.6 Machined ASTM E8 tensile specimen
79
4.4 Differential scanning calorimetry
The differential scanning calorimetry (DSC) was conducted with the TA
Instrument SDT 650 in this investigation at OSU in order to evaluate the (α + β) and β transus. The specimen was placed in an alumina crucible with an empty alumina crucible as reference. The specimen chamber is evacuated and constantly flowed with argon gas.
The empty alumina crucible was tested before the sample to obtain the baseline, which was deducted for the results of the sample. The DSC instrument, namely the thermocouple, has been calibrated regularly. The calibration process includes both the temperature and sensitivity calibrations, which are specific to the crucible, the heating rate, and the protective gas. The temperature calibration eliminates the deviation between the temperature measured by the thermocouple and the actual temperature of the sample.
Sensitivity correction ensures the accurate conversion between thermocouple signal and heat flow power. The chamber temperature was increased from room temperature to
1100°C at the rate of 20°C/min.
4.5 Mechanical property test
The machined ASTM E8 specimen from as-cast ingots was shown in Figure 4.6.
The gauge length is 25.4 mm and the gauge diameter is 4 mm. The tensile tests were
conducted with MTS tensile frame (MTS Criterion Model 32) at room temperature with
the displacement rate controlled at 0.004 mm/s (steady state). Three tensile tests were
80
conducted. The average tensile properties and representative true stress-strain curves were reported.
4.6 Results and discussion
4.6.1 Microstructure characterization
Figure 4.7 exhibit the SEM images of as-cast T65-0.05BC specimens with different magnifications. It can be seen that in the as-cast microstructure, the intragranular α phases are fine platelets but in neither basketweave nor lamellar forms.
The grain boundaries are also free from continuous α platelets and precipitation-free zone
(PFZ). The morphology on grain boundaries are discrete ellipsoidal α phases and small intermetallic particles, which are likely TiB since boron has very limited solubility in either α or β phases. It should be noted that in between the primary α platelets, the ultra- fine α precipitates are observed with both SEM and TEM in Figure 4.7 and Figure 4.8(a) and (b), assumed to be secondary α precipitation from β matrix. Their morphology is similar to the duplex microstructure in α-β titanium alloys, except they are smaller (with average width of 36 nm). The formation of secondary α phase in the as-cast status is usually unexpected in conventional titanium alloys, which are commonly achieved by low temperature heat treatment or hot working [1.1]. The hypothetical cause for the formation of this fine secondary α phase is the strong iron partitioning in titanium alloy shown in Figure 2.2. The high-temperature saturated post-primary-α β matrix becomes super-saturated at low temperatures during cooling, leading to the driving force for
81
precipitation of the secondary α, and the fast diffusivity of iron ensures that this process
is kinetically possible. Based on this hypothesis, this process should be diffusion-
controlled and the secondary-α-β interface should have a strong iron concentration
gradient. The HAADF-STEM image and correlated SETM-EDS scans of the as-cast
microstructure in Figure 4.8 partially confirms the strong partitioning of iron in between
α and β phases in this alloy. It can be seen that compared with aluminum, the partitioning effect of iron is very strong, resulting in a close-zero concentration in α phase.
Additionally, the aggregation of iron content at the α-β interface also indicates that a fast kinetic process triggers the as-cast phase transformation. Since the secondary α phase is
too fine for the EDS resolution, atomic probe tomography is preferred to further
investigate and confirm the partitioning in secondary α phase formation.
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Figure 4.7 SEM images of as-cast Ti-6Al-5Fe-0.05B-0.05C under different magnifications
83
Figure 4.8 STEM characterization of the microstructure of Ti-6Al-5Fe-0.05B-0.05C: (a)
Bright field STEM image with a inserted selected area diffraction pattern showing the
Burgers orientation relationship between α and β phase, (b) HAADF-STEM image, (c)
STEM EDS maps, and (d) STEM EDS line.
84
4.6.2 As-cast tensile property
The representative as-cast true stress-strain curve of T65-0.05BC is shown in
Figure 4.9, while Table 4.1 compares the tensile properties with three other as-cast α-β titanium alloys. The reported mechanical properties of T65-0.05BC in Table 4.1 are mean values with standard deviations from all three tests. The results show that the as-cast T65-
0.05BC possesses relatively higher yield and ultimate tensile strengths than as-cast Ti-
6Al-4V, Ti-6Al-4V ELI, and Ti-5Al-2.5Fe alloys. This increase in strength can be attributed to the fine α platelets and ultra-fine secondary α precipitates as shown in Figure
4.7 and Figure 4.8. It should also be noted that the oxygen content in this alloy, 0.061 wt.%, is considered lower than common commercial titanium alloys [1.1]. Since oxygen is generally regarded to have a strong hardening effect, this alloy is capable of reaching higher strength with typical commercial oxygen level (0.2 wt.%) [1.1]. However, the formation of this secondary α during cooling is uncommon in α-β titanium alloys, where it is usually achieved via the combination of an intermediate solution heat treatment followed by a lower temperature aging treatment [1.1]. Different heat treatments can be applied to tailor the fractions of primary and secondary α and β phases in order to adjust the strength-ductility balance for specific applications in this alloy. A β-annealing attempt is conducted in this investigation. It should be mentioned that this alloy can not only be used for casting applications, but also for powder production and subsequently powder metallurgy for additive manufacturing since its β→α phase transformation can occur under even fast cooling condition.
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Figure 4.9 ASTM E8 tensile result of as-cast Ti-6Al-5Fe-0.05B-0.05C
Table 4.1 Tensile properties of as-cast Ti-6Al-5Fe-0.05B-0.05C in comparison with as- cast Ti-6Al-4V [4.16], Ti-6Al-4V-ELI [4.16] and Ti-5Al-2.5Fe [4.17, 4.18].
As-cast Yield Strength Ultimate Tensile Strength Elongation Material (MPa) (MPa) (%) Ti-6Al-5Fe- 1023 20 1136 40 3.71 0.13 0.05B-0.05C Ti-6Al-4V 895 1000 8 Ti-6Al-4V ELI 825 896 10.5 Ti-5Al-2.5Fe 820 900 6.0 86
4.6.3 (α + β) and β transus
As shown in Figure 4.10, the onset and offset points of the β transus peak are
marked. The onset and offset temperatures should correspond with the (α + β)
transformation temperature and the β transus respectively. The β transus is close to
CALPHAD prediction, while the (α + β) transformation temperature is higher than the
calculation. This could potentially be due to the thermodynamic description of TiFe
phase in PanTi database. As can be seen in Figure 4.1, there is a (α + β + TiFe) phase
region below the (α + β) phase region. This region exists in the isopleth because the TiFe
phase in PanTi database is described as an independent intermetallic phase. However, in
recent thermodynamic assessment [4.19], TiFe phase should be regarded as an ordered
BCC_B2 phase described together with β-Ti, which should eliminate this 3-phase region,
and further raise the calculated (α + β) transformation temperature.
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Figure 4.10 DSC signal curve of T65-0.05BC through heating. The dashed lines are applied to determine the onset and offset temperatures.
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4.7 Conclusion
In summary, the preliminary results have shown a new cost-effective cast titanium
alloy Ti-6Al-5Fe-0.05B-0.05C. The substitution of Fe for V significantly reduces the raw
material cost, and the Ti-Al-Fe alloy system has characteristics for good castability
according to CALPHAD calculations. The micro-tension test results show that this new
alloy has very high strength due to the primary α lamellae and ultra-fine secondary α
precipitates. The strong partitioning of Fe to the -phase was demonstrated. It is believed
that the strength-ductility balance of the new alloy can be tailored by heat treatment and thermomechanical processing, adjusting the fractions of α (primary or secondary) and β phases, for various applications. Better manufacture process shall be further investigated to acquire improved as-fabricated microstructure and properties, and will also be beneficial for further processing.
4.8 References
[4.1] F.H. Froes, D. Eylon (eds.), Titanium Net Shape Technologies, TMS,
Warrendale, 1984.
[4.2] D. Eylon, F.H. Froes, Titanium Casting - A Review, Titanium Net Shape
Technologies, in: F.H. Froes, D. Eylon (Eds.), Titanium Net Shape
Technologies, TMS, Warrendale, 1984, pp. 155-178.
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[4.3] Y.G. Li, X.D. Zhang, P.A. Blenkinsop, N. Saunders, N.A. Walker, P.D. Spilling,
C. Small, in Titanium `95: Science and Technology, the Proceedings of the 8th
World Conference on Titanium, ed. P. A. Blenkinsop, E. J. Evans and H. M.
Flower. Institute of Materials, 1996, pp. 2317–2324.
[4.4] Y.G. Li, P.A. Blenkinsop, M.H. Loretto, N.A. Walker, Effect of alul11iniul11on
ordering of highly stabilised beta-Ti-V-Cr alloys, Mater. Sci. Technol., 14
(1998) 732-737.
[4.5] Y.G. Li, P.A. Blenkinsop, M.H. Loretto, N.A. Walker, Structure and stability of
precipitates in 500°C exposed Ti–25V–15Cr–xAl alloys, Acta Mater, 46 (1998)
5777-5794.
[4.6] K.H. Borowy, K.H. Kramer, in: G. Lütjering et al. (Eds.), Titanium Science and
Technology, Deutsche Gesellschaft für Metallkunde E.V., 1985, pp. 1381-1386.
[4.7] P.J. Bania, A.J. Hutt, R.E. Adams, W.M. Parris, in: F.H. Froes, I.L. Caplan
(Eds.), Titanium ‘92 Science and Technology, TMS, Warrendale, 1993, pp.
2787-2794.
[4.8] H. Fujii, K. Takahashi, Nippon Steel Tech. Rep,. 85 (2002) 113-117.
[4.9] D.J. Lin, J.H.C. Lin, C.P. Ju, Structure and properties of Ti–7.5Mo–xFe alloys,
Biomaterials, 23 (2002) 1723-1730.
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[4.10] H.C. Hsu, S.K. Hsu, S.C. Wu, C.J. Lee, W.F. Ho, Structure and mechanical
properties of as-cast Ti-5Nb-xFe alloys, Materials Characterization, 61 (2010)
851-858.
[4.11] S. Roy, S. Suwas, S. Tamirisakandala, D.B. Miracle, R. Srinivasan,
Development of solidification microstructure in boron-modified alloy Ti–6Al–
4V–0.1B, Acta Mater, 59 (2011) 5494-5510.
[4.12] S. Tamirisakandala, R.B. Bhat, J.S. Tiley, D.B. Miracle, Processing,
microstructure, and properties of β titanium alloys modified with boron, J.
Mater. Eng. Perform., 14 (2005) 741-746.
[4.13] R. Banoth, R. Sarkar, A. Bhattacharjee, T.K. Nandy, G.V.S. Nageswara Rao,
Effect of boron and carbon addition on microstructure and mechanical properties
of metastable beta titanium alloys, Mater. Des., 67 (2015) 50-63.
[4.14] R. Sarkar, P. Ghosal, T.K. Nandy, K.K. Ray, Structure–property correlation of a
boron and carbon modified as cast β titanium alloy, Philos. Mag., 93 (2013)
1936-1957.
[4.15] F.A. Crossley, Grain refinement of titanium alloys, U.S. Patent No. 4420460,
1983.
[4.16] E.E. Billinghurst, Jr., NASA Tech. Pap., 3288 (1992) 9-10.
91
[4.17] R. Boyer, G. Welsch, E.W. Collings (eds.), Materials Properties Handbook:
Titanium Alloys, ASM International, Materials Park, 1994.
[4.18] J. Black, G. Hastings (eds.), Handbook of Biomaterial Properties, Chapman &
Hall: London, New York, 1998.
[4.19] H. Bo, J. Wang, L. Duarte, C. Leinenbach, L. Liu, H. Liu, Z. Jin,
Thermodynamic re-assessment of Fe-Ti binary system, Trans. Nonferrous Met.
Soc. China, 22 (2012) 2204-2211.
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Chapter 5 . Manufacture of Lab-scale Prototype Casting Automotive Connecting Rod
5.1 Introduction
In this investigation, the new low-cost casting titanium alloy developed in Chapter
4 will be conducted with casting manufacture process, which is a low-cost manufacture process preferred by civil application, such as automotive industries. Since civil application requires massive production, the current precision casting techniques mentioned in Chapter 2 do not meet the industrial cost consideration and must be modified for lower cost. The specific application, the automotive engine connecting rod, is required by the industrial partner of this project, General Motors. And the permanent mold casting technique is applied as the low cost casting option and optimized accordingly to meet the requirement for titanium casting.
In this chapter, a lab-scale casting manufacture framework is illustrated, including equipment and mold design, OSU facility setup, casting simulation, melting and casting operations, and the ballpark manufacture cost analysis.
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5.2 Lab-scale manufacture casting framework setup
This section will illustrate the concept of this lab-scale casting framework, the melting/casting equipment design, the permanent mold concept and design, and equipment installation in OSU facility.
5.2.1 Basic concept of the framework
As mentioned in Chapter 1, this investigation is motivated by the idea of low cost alloy production, so the design of casting manufacture process is also cost-directed. The casting process can be generally considered as two steps: melting and casting.
Specifically, for titanium alloys, as mentioned in section 2.2.1, there are not many choices for the melting technique since the molten titanium is at high temperature and has high reactivity, and thus requires to be processed in vacuum or inert atmosphere with
CHM setup. In this project, for the convenience of setup in OSU facility, the small ISM is chosen as the melting option. As for the casting techniques, the current mature ones, rammed graphite casting and investment casting, both have the cost disadvantages.
Although these two casting methods provide more efficient and economical ways of producing large volumes of titanium components, there are still problems remaining for these two methods. Both conventional rammed graphite mold casting and investment casting definitely can produce very complex shape, yet the entire procedure is complex compared with permanent mold casting for massive production. The other significant impedance in these two techniques is the non-permanent mold setup. In both techniques, each cycle of casting production requires a repeated molding process because the mold
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cannot be reused, and it significantly increases the cost of casting operation, raising the
energy and labor cost. Therefore, in the concept of this framework, the permanent metallic mold casting is determined in order to address the cost issue in molding, which is the most important concept in this investigation.
There are a few challenges in applying the permanent metallic mold casting to titanium alloys, mainly due to the high reactivity and viscosity of the titanium melt. For high reactivity, lowering the melting temperature is an efficient way to decrease the potential interaction between the molten metal and metallic mold, which is introduced in
Chapter 2 and addressed during alloy design in Chapter 4. As for viscosity, in general, the viscosity of titanium alloys is roughly five times of the one of aluminum and magnesium alloy [5.1, 5.2]. The viscosity is limited by the nature of the solvent element, and cannot be modified in numbers of magnitude by alloy design. The only possibility is to raise the melt temperature, which is dependent on the melting and molten metal transfer techniques. This topic is not within the scope of this project and therefore not addressed.
5.2.2 Equipment and mold design
The equipment concern in this framework is mainly about the choice of melting
technique since the casting/molding technique is already determined to be permanent
metallic mold casting. For the current melting techniques listed in Chapter 2, multiple
manufacturers were consulted and ISM was determined as the finalized option. This is
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mainly due to the budget and facility spatial limitation on OSU site. The other two CHM techniques, PCHM and EBCHM, are beyond the budget of the project (~$1.6m) and
cannot be spatially installed in OSU foundry. ISM, on the contrary, is small in size and sufficient in melting lab-scale volume to produce conceptual prototype parts for laboratory validation purpose.
The permanent metallic mold casting in this project was modified into a high- temperature ceramic-coated permanent mold casting. This design is inherited from the
ceramic mold concept from the investment casting. The high temperature ceramic mold
in the investment casting is used for multiple reasons, including low reactivity with the
molten titanium, and low heat transfer rate enhancing filling fluidity. Therefore, in this
project, instead of ceramic molding, a thin layer of ceramic coating is applied with air
spraying on the permanent metallic mold to prevent the reaction between the mold and
molten metal. The detailed design applied common H13 tool steel as the mold and
zirconia (ZrO2) as the coating material. A number of studies were carried in the
investment casting field to investigate different ceramic compounds as molding materials
[5.3-5.6]. Based on the industry need for molten metal reactivity and cost compensation, different coating materials can be applied in the industry-scale production. The pre- casting and post-casting mold pictures are shown in Figure 5.1. As can be seen, the mold was not damaged from molten metal contact, but the coating was peeled off from the mold surface. This is potentially due to the poor adhesion from the air spraying technique. The common coating spraying technique in industry is thermal spraying,
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which should provide better adhesion with mold surface and allow the ceramic coating to endure multiple casting cycles. The mold geometries are two-halves of a connecting rod and casted separately.
5.2.3 OSU facility setup
The OSU facility lab-scale manufacture framework setup is referred to the installation of ISM melting-casting furnace. The ISM melting-casting furnace was designed and manufactured by PVT Inc., an InductoTherm Group company, in which several key concepts were determined by the OSU team to meet the requirements in the project and installation needs in OSU facility. The overall picture of this unit and several details are shown in Figure 5.2. The main features of this unit compose of a 500g capacity (referring to titanium alloys) ISM system, a resistance mold heating unit, and a mechanical tilting system. The first two features are both located inside a vacuum chamber capable of reaching high vacuum for titanium melting, and the last one allows the ISM system to be tilted to gravity-pour the molten metal into the mold.
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Figure 5.1 (left column) Pre-casting mold with ZrO2 coating of (top to bottom) top half and bottom half, and (right column) post-casting mold pictures of (top to bottom) top half, bottom half, and core
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DI water‐cooled copper crucible
Funnel & Mold Location
Resistance mold heater
Figure 5.2 (top) OSU facility ISM-Casting furnace and (bottom) layout in the vacuum chamber 99
A few features were considered in this melting/casting setup in ahead, but
abandoned due to the budget limitation. These features are listed below and can be
considered for further improvement or industry-scale production:
• Mold cooling system: allowing the mold to be cooled faster after the filling and
leading to better microstructure/final component properties
• Centrifugal filling-assisted system: allowing higher filling driving force and thus
faster/better filling performance
• Immersive pyrometer: in-time monitoring of melt temperature
• Computer-controlled tilt-pour system: allowing precise control of the tilt-pour in
the casting operation, and enhancing the stable repeatability of the process
5.3 Casting simulation
In this section, the casting simulation was conducted as an effort to introduce
computational method as guidance for manufacturing optimization. Based on the mold
design and actual casting setups in the previous section, the casting model was
constructed and simulated with the commercial casting simulation code, EKK
CAPCAST, in order to predict the casting performance of T65-0.05BC alloy in the metallic permanent mold casting, including cavity filling, solidification sequences, and temperature profiles. The selected results are shown in Figure 5.3. The filling velocity, 100
solidification, and flow simulation shall reveal any potential inhomogeneity in molten
metal flow and hot spot, which are potential sources for gas entrapment and shrinkage defects. These imperfections cannot be completely eliminated but can be maximally avoided by reducing the gradients around the high intensity locations in the contours. Due to the lack of parameters for the new alloy and mold coating effect, the simulation results might not reflect some details in this setup. However, the trend of filling and solidification is still useful in improving mold and process design.
5.4 Prototype casting in OSU facility
In this section, the details of the melting-casting process will be illustrated, including parameters for power, pouring and post-casting operations.
5.4.1 Raw materials
In order to ensure that the process is similar to industrial process, it is necessary to use commercial pure grade raw materials instead of high purity ones used for alloy design in Chapter 3. In this process, the raw materials are CP-2 titanium (99.9%, ATI Metals),
CP aluminum (99.9%, Alfa Aesar), CP iron (99.9%, Alfa Aesar), TiB2 (99.5%, Alfa
Aesar), carbon slug (99.5%, Alfa Aesar). In the recommendation for further
industrialization/commercialization of the process, the CP iron and carbon slug shall be
replaced with certain recycled steel scraps.
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Figure 5.3 EKK simulation of (top row) final temperature/solidification and (bottom row) filling time
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5.4.2 Melting operation
In this process, the raw materials were stacked inside the DI water cooled copper
crucible and melted in the induction field. The stacking of raw materials is designed as
shown in Figure 5.4 in order to reach maximal coupling between induction field and
materials. The pure titanium bars, the highest melting point charge material, accounted
for most mass of the alloy and was placed starting from the bottom of the crucible to
make sure that it was located at the strongest position in the induction field. The
aluminum buttons and iron granules were placed on the top of titanium to ensure that
they fell into the molten titanium but not reach the bottom of the crucible first and
instantly solidified. As for TiB2 and carbon slug, due to their trace amounts and powder-
nature, they were placed among titanium bars to prevent from splashing out of the
crucible by induction stirring.
The melting operation started with the pre-heating of the raw materials to avoid
thermal explosion due to partial heating of the raw materials. The thermal explosion can
cause safety issues in the operation and also lead to splashing of the molten metal, resulting in inaccuracy of the alloy composition. The power of the induction power output was set at 10 kW (20% of maximal power) for 15-20 minutes. From the observation port on the chamber, the pre-heating should be done once the heat radiation
(glowing) of the charge materials was well distributed. Then the power should be gradually increased to maximal level, which is 50 kW for OSU equipment. As soon as the charge materials reached full molten state, it should be kept at the molten state for 10-
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15 minutes to ensure that the charge materials were completely melted. It must be noted that during the melting and casting process, the chamber should be vacuumed and back- filled with Ar to approximately 80 torr (roughly 1/10 of atmospheric pressure) to prevent aluminum evaporation during melting. The evaporation of aluminum could severely affect the final alloy composition.
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Crucible
CP‐2 Ti bars
TiB2 powder and C slug CP‐Fe granules CP‐Al buttons
Figure 5.4 Schematics of charge material stacking in DI-water-cooled copper crucible
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5.4.3 Casting operation
After the molten metal is stable and considered fully melted, the tilt-pour gravity casting is conducted with the preheated mold. In this investigation, the H13 metallic mold is preheated to 700°C to increase the molten metal fluidity as much as possible. Before tilting, the induction power is turned down to 25 kW (50% of max power) to prevent the molten metal splashing while keep the superheat as long as possible. Then the tilt needs to happen instantaneously to conduct the casting before the superheat dissipates. Since this tilt process in OSU facility is not computer-controlled, several “trial-and-error” practices are needed to acquire good casting. This process should have increasing repeatability with the computer-controlled system as mentioned in sub-section 5.2.3.
After the filling is finished, the mold heat power is turned off to let the mold cool down radiantly.
5.5 Casting results and post-assessments
The prototype casting is shown in Figure 5.5. Overall, the permanent metallic mold casting concept is successfully proven. The metallic mold is successfully filled under gravity-pour condition, which is the least-driving-force cavity filling condition.
And there are no significant unfilled or shrinkage portions, indicating that this process can be commercialized with even better performance.
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Several casting defects are observed on the component surface, including gas porosity
and surface roughness. These defects are due to the experimental measures applied in this project, and can be resolved by using mature techniques in industry setup. It should also be mentioned that the mold pre-heat temperature could be lower with the filling-assisted technique.
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Figure 5.5 Prototype casting T65-0.05BC connecting rod
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5.6 Manufacture cost analysis
The cost analysis includes three aspects: raw materials, energy and labor in
operations, and equipment. Only the first two aspects are counted in the final cost
analysis comparison. The preliminary equipment cost is listed as references and not
accounted in the overall analysis. The actual equipment cost analysis depends on the
detailed amortization plan by industry. Since the post-treatment of the new alloy is still
under investigation, this cost analysis does not include post-treatment estimation but only
as-fabricated.
5.6.1 Raw Materials
When estimating raw materials, the cheapest options are considered first to acquire the lowest potential cost. In this analysis, the raw materials cost are determined based on the available market prices of pure element, master alloy, and recycled alloy.
The raw materials analysis is shown in Table 5.1. As can be seen, the new titanium casting alloy is strongly competitive in raw materials cost with either conventional PM
stainless steel or Ti-64. However, it should be noted that the available market price for
P/M steel powder is based on the open import price from India, and the price in the
United States might vary depending on the suppliers. The effect of the actual P/M steel powder price is shown in Figure 5.6 after the final cost analysis summary.
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Table 5.1 Cost analysis of raw materials
Cast T65-0.05BC Est. Market Required Cost per Cost per part Price weight weight ($) ($/kg) (kg) ($/lb) CP-2 Ti1 3.747 0.45 1.68615 CP-Al2 2.45 0.0304 0.07448 Recycled high 0.209 0.0255 0.0053295 carbon steel3 4 TiB2 powder 50 0.000815 0.04075 Total 0.506715 1.8067095 1.621 Total for 100,000 parts ($) 180,670.95 Cast Ti-64 CP-2 Ti1 3.747 0.45 1.68615 CP-Al2 2.45 0.0304 0.07448 CP-V4 29.98 0.020269 0.607664 Total 0.506716 2.3909488 2.144781 Total for 100,000 parts ($) 239,094.88 P/M steel Stainless steel 3.263 0.83 2.70829 1.48 powder5 Total for 100,000 parts ($) 270,829.00 Sources for market prices: 1Argus Metals International Global non-ferrous market prices, news and analysis, Jan 30, 2018 2MetalMiner 3Rockaway Recycling 4 Alibaba, TiB2 powder price varying from $20-100 per kg. A mid-value is taken for estimation; CP-V price varying from $20-40 per kg. A mid-value is taken for estimation; It should be noted that raw material prices are generally lower in Chinese market, and US providers’ price might change the estimation. 5Commodity lowest available import price from India (prices at 5 tons basis, largest quantity found). Price in US might vary.
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5.6.2 Energy and Labor
The energy and labor costs are estimated based on the available specifications of the manufacture equipment (power, size, etc.), the energy cost rate (state-based, use
Michigan state in this analysis), and the hourly compensation costs in manufacturing in
the United States. The comparison is shown in and Table 5.3. It can be seen that the
conventional P/M process is labor-saving compared with the current estimation on ISM
casting due to the differences in the number of components per operation. In the final cost
analysis, the ISM casting process cost is applied for both T65-0.05BC and Ti-64.
However, it should be noted that this estimation can vary depending on the actual casting
and P/M setups in industry with different production setup.
5.6.3 Equipment
The equipment cost is based on open sources and confidential discussions with
manufacturers listed in and Table 5.3 with energy and labor cost. The equipment cost of
ISM furnace is high due to the specific requirements from the processing of titanium
alloy.
5.6.4 Final Cost Analysis Comparison
As can be seen from the final cost comparison in Table 5.4, the starting target is
closely achieved with 38.7% weight reduction at a cost of less than $2.63 per pound of
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weight saved for a finished titanium connecting rod casting compared with the baseline part of P/M steel. In addition, it should be noted that with the same casting process, the new designed Ti alloy is also very competitive in raw material price with Ti-64 with the same part weight. Although the DOE project target of a cost of less than $2.25 per pound of weight saved is not reached, the weight reduction has exceeded the DOE project target
(30%). It should be noted that the actual cost analysis might vary from the ballpark analysis in this project depending on the equipment manufacture and raw materials providers.
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Table 5.2 Equipment, energy and labor cost analysis for cast T65-0.05BC
Cast T65-0.05BC Equipment cost Equipment estimate Additional Information ($) 10kg capacity with centrifugal caster ISM 2,750,000 Power: 400 kW Mold 30,000 Account for top and bottom molds Energy Operation time Energy cost Parts per Processes consumption (hr/operation) ($/operation) operation (kWh/operation) Melting + 0.2 80 8.84 20 Casting Total energy cost for 100,000 parts 44,200 ($) Mold heat up and Cost for 1.0 Cost per vacuuming 100,000 operation ($) Melting + Casting 0.3 parts ($) Labor cost Mold cooling 0.2 (hr/step) Mold 0.3 70.285 351,270 cleaning/spraying Total 1.8 Total add-up for 100,000 parts ($) 395,470
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Table 5.3 Equipment, energy and labor cost analysis for P/M steel
P/M steel Equipment cost Equipment estimate Additional Information ($) Compact hydraulic 350,000 Power: 90 kW Working area: 1.6x2.0m press Sintering Power: 250 kW Chamber size: 300,000 furnace 0.63x0.125x10.5m Mold 30,000 Account for top and bottom molds Energy Operation time Energy cost Parts per Processes consumption (hr/operation) ($/operation) operation (kWh/operation) Cold 0.33 30 3.315 100 Compacting Sintering 1.5 400 44.2 100 Total energy cost for 100,000 parts 47,515 ($) Compacting and 0.33 Cost for mold cleaning Cost per Labor cost 100,000 operation ($) (hr/step) parts ($) Sintering 1.5 71.4249 71424.9 Total add-up for 100,000 parts ($) 118,906.75
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Table 5.4 Cost analysis summary
Cast T65-0.05BC Cast Ti-64 P/M stainless steel Raw materials ($) 180,670.95 239,094.88 269,010.00 Labor cost ($) 351,270 351,270 71,424.9 Process energy cost ($) 44,200 44,200 47,515 Total ($) 576,140.95 634564.88 387,916.75 Total per part ($) 5.76 6.35 3.88 Part weight (lb) 1.12 1.12 1.83 Total per lb ($) 5.16 5.67 2.12 Weight saved per part Negligible 38.7% compared with (%) due to small Cost increase density 2.63 compared with ($/lb) difference
8
6 saved
4 weight
of
2 pound
per 0 $
12 10 8 6 4 2 0 in
Current ‐2 price The target setup estimation: in project: ‐$2.63 ‐$2.25 ‐4 Increase
‐6 Stainless steel powder price ($ per kg)
Figure 5.6 The effect of P/M stainless steel powder price on final cost estimation
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5.7 References
[5.1] T.W. Chapman, The Viscosity of Liquid Metals, A.I.Ch.E. Journal, 12 (1966)
395-400
[5.2] T. Ishikawa, P. Paradis, J.T. Okada, Y. Watanabe, Viscosity measurements of
molten refractory metals using an electrostatic levitator, Meas. Sci. Technol., 23
(2012) 025305
[5.3] H. Jiang, B. Li, J. Guo, S. Dong, Z. Li, China Foundry. 4 (1998) 22-24.
[5.4] Jiafang Wang, Jiannong Wang, J. Yie, Spcial Casting & Nonferrous Alloy. 5
(2002) 40-42.
[5.5] B. Li, J. Shang, Z. Guo, J. Aeronautical Mater, 21 (2001) 32-35.
[5.6] Y. Li, Master Dissertation, HIT. 2006.
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Chapter 6 . Conclusions and Future Perspectives
6.1 Conclusions
In this 4-year project/5-year PhD investigation, we started from exploring new low-cost titanium alloy systems with the guidance of Integrated Computational Materials
Engineering – Calculation of Phase Diagram (ICME-CALPHAD) framework, and determined several promising alloy systems, alloying elements including Ti-Al-Fe-Mn-
B-C, that can allow further development (lab-scale and industry-scale) in the future.
During the new alloy system exploration, considerable amount of scientific work was completed regarding CALPHAD modeling, microstructure morphology, mechanical properties, and processing of the new alloy systems. The relevant ternary titanium alloy systems, Ti-Al-Mn and Ti-Al-Fe, were examined and reassessed both computationally and experimentally, leading to more comprehensive understanding of these alloy systems.
Based on this above study, a new low-cost casting titanium alloy was developed,
Ti-6Al-5Fe-0.05B-0.05C (designated as T65-0.05BC), which has competitive as-cast strength level with current commercial titanium alloys but relatively low ductility.
However, in this project, since the target application, automotive engine connecting rod,
117
is a fatigue-limited component and does not require high ductility, this alloy shall suffice as the deliverable alloy in this project.
After designing and confirming the potent of the new alloy, we designed and built
a lab-scale manufacturing equipment aided by computational tools at OSU labs to
conceptually prove that this alloy can be manufactured with permanent mold casting,
which is low-cost and potentially a mass production process. We produced a lab-scale
prototype connecting rod demonstrating the capability of permanent mold casting of this
new alloy. Based on the results from lab-scale manufacturing, several recommendations
are made for commercialization/industrialization of this process.
Also, in order to address that this alloy and the related manufacturing process we
proposed are cost-competitive with the current alloy and manufacturing techniques, a cost
analysis was conducted in this investigation. The results showed that compared with current commercial titanium alloy, Ti-6Al-4V, and conventional powder metallurgy stainless steel, this new casting titanium alloy has the potential to replace them in actual manufacturing environment.
In addition to the major achievements listed above, there are a few issues that need to be investigated about this new alloy and the manufacturing process. For example, the heat treatment of this new alloy needs to be further investigated to control the stable microstructure. The permanent mold casting process requires more precise design to reach maximum efficiency during operation.
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In summary, this investigation made significant progress in applying
computational tools to design new titanium alloys and manufacturing processes in an
accelerated rate. Evidently, the deliverable new low-cost casting titanium alloy has been
designed with a lab-scale validated manufacturing framework. This lab-scale framework
can be extended to industry-scale and beneficial to the application of titanium alloys in
civil applications, which in this case are automotive components. Besides the effort done
in this project, more investigations are strongly suggested to continue and further develop
this new alloy/alloy system with this manufacturing process (or other low-cost
manufacture options).
6.2 Ongoing investigations and future perspectives
6.2.1 Further investigation of thermodynamic and kinetics of Ti-Al-Fe system
As briefly mentioned in Chapter 2, the thermodynamics and kinetics descriptions
of Ti-Al-Fe system still require further work to be done to construct a more accurate
database for further alloy design. OSU team is currently constructing a customized Ti-Al-
Fe database. The main difficulty is to compromise among the assessed binary systems
about the thermodynamic definitions in the database programming. The experimental Ti-
Al-Fe phase diagram is briefly summarized as shown in Figure 6.1 [6.1]. The Ti-Al
binary system has no conflicts with the other two binary systems since most phases are compounds and no ordered solid solution phases. The main existing problem is the phase definition of ordered BCC phase in Ti-Fe and Al-Fe system. In the latest Ti-Fe binary 119
system description, TiFe, an intermetallic compound, is regarded as an ordered BCC
solid solution phase with 2 sub-lattices (3 when counting vacancy) [6.2]. However, in Al-
Fe binary system, there exists two ordered BCC solid solution phases, and they are
modeled with 2 and 4 sub-lattices respectively. However, due to their internal modelling conflicts, one must be forbidden in order to show the other when calculating phase diagrams. But in the experimental phase diagram, all ordered BCC solid solution phases
are shown. This definition conflict issue needs to be addressed to construct the Ti-Al-Fe
thermodynamic database.
A few preliminary experiments were conducted to systematically investigate the
effect of iron in Ti-Al-Fe system. Specimens with compositions Ti-6Al-xFe were casted
with arc melter. The specimens were firstly β-annealed at 1100°C/1hr and (α + β) aged at
800°C/1hr. The aged microstructures are shown in Figure 6.2. As can be seen, the phase
evolution with increased iron content is in some degree consistent with the calculated result in Figure 4.1. There appears to be minor formation of intermetallic phases when the
iron content reaches 12 wt.%, which is suspected to be primary TiFe precipitating from β
matrix. The microstructures become less homogenous and wider PFZ as iron content
increases, which is possible caused by the increased anisotropy of matrix-β lattice.
120
Figure 6.1 Experimental 1000°C isothermal section of Ti-Al-Fe system [6.1].
121
6.2.2 Further development of the new alloy, T65-0.05BC
One issue about the current new alloy is the design of heat treatment procedure.
Usually, the as-cast component requires heat treatment, including β hot-isostatic-pressing and α-β aging, to acquire a defect-free final microstructure. A few preliminary time- temperature-transformation (TTT) experiments were conducted as an attempt to investigate the phase transformation kinetics. However, an existing problem in the current new alloy is that the β→α phase transformation rate is too fast to be precisely controlled (within a few minutes) potentially due to the fast iron diffusion. As shown in
Figure 6.3, specimens all underwent β-homogenization at 1100°C/1hr and different aging treatments. Unfortunately, since the phase transformation process is too fast, the kinetic information cannot be extracted from these experiments. For example, for 800°C aging treatment, there is no measurable α phase volume fraction among 15, 30, and 90 minutes, indicating that the phase transformation is complete within 15 minutes. For heat treatment below 15 minutes, it is difficult to obtain reliable isothermal temperature profiles. OSU team is currently exploring possible methods for this type of heat treatment requirements.
122
Ti‐6Al‐3Fe Ti‐6Al‐4Fe Ti‐6Al‐5Fe
Ti‐6Al‐6Fe Ti‐6Al‐8Fe Ti‐6Al‐10Fe
Ti‐6Al‐12Fe Ti‐6Al‐14Fe
Figure 6.2 SEM images of heat treated Ti-6Al-xFe: 1100°C/1hr + 800°C/1hr
123
15 min 30 min 60 min
15 min 30 min 90 min
5 min 15 min
Figure 6.3 TTT experiments of Ti-6Al-5Fe-0.05B-0.05C: β-homogenization at
1100°C/1hr + (row 1) 750°C, (row 2) 800°C, (row 3) 850°C. The aging time is tagged.
124
6.2.3 Involvement of Mn in alloy design
As mentioned in Chapter 4, manganese is temporarily excluded in this
investigation, but OSU team is trying to produce several T65-0.05BC specimens with different manganese contents for evaluation to see if manganese is valuable to be
introduced in future development. A few preliminary experiments were conducted to
investigate potential manganese effect on microstructures as shown in Figure 6.4. As the
Fe-to-Mn ratio decreases, the microstructure becomes less anisotropic and more typical
basket-weave form, indicating that iron dominates the primary α phase morphology.
6.2.4 Mold coating investigation
In the original plan, the interaction in the mold-coating-casting is supposed to be
quantitatively investigated, and there are a few results generated in the previous quarterly
reports. However, due to the adhesion issue mentioned in Chapter 5, this investigation is
postponed until more stable/adhesive coating quality can be achieved. Several molten
titanium-ceramic couple experiments were conducted to investigate the interfacial
reaction as shown Figure 6.5. This experiment was conducted inside the plasma arc
melter using the plasma arc to melt the alloy button and allowing it to drop on the
ceramic (ZrO2) button to imitate the interfacial reaction at the coating-mold contact
locations. This experimental setup can be improved with more realistic setup, which better imitates the actual coating conditions.
125
Ti‐6Al‐5Fe Ti‐6Al‐4Fe‐1Mn Ti‐6Al‐3Fe‐2Mn
Ti‐6Al‐2Fe‐3Mn Ti‐6Al‐2.5Fe‐2.5Mn Ti‐6Al‐1Fe‐4Mn
Ti‐6Al‐5Mn
Figure 6.4 SEM images of Ti-6Al-xFe-yMn: β-homogenization at 1100°C/1hr +
800°C/1hr
126
(a)
(b) (c)
Figure 6.5 (a) The molten titanium – ceramic couple experiment setup in plasma arc
melter, and SEM images of the ZrO2-titanium interfaces of (b) Ti-6Al-2Fe-2Mn and (c)
Ti-6Al-1Fe-3Mn
127
6.3 References
[6.1] M. Palm, J. Lacaze, Assessment of the Al-Fe-Ti system, Intermetallics, 14
(2006) 1291-1303.
[6.2] B. Sundman, I. Ohnuma, N. Dupin, U.R. Kattner, S.G. Fries, An assessment of
the entire Al-Fe system including D03 ordering, Acta Mater, 57 (2009) 2896-
2908.
128
Chapter 7 Supplementary Results
0.1BC 0.2BC 0.3BC
Figure 7.1 SEM images of as-cast Ti-6Al-5Fe-xB-xC: (top) low magnification, and
(bottom) high magnification
129
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139