DISPERSION STRENGTHENED TITANIUM ALLOYS
PRODUCED VIA DIRECT LASER DEPOSITION
A Thesis
Presented in Partial Fullfillmt!nt of the Requirements for
the Degree of Maser of Science in the
Graduate School of The Ohio State University
By
Craig Alan Brice, B.S.
*****
The Ohio State University
2000
Master's Examination Committee:
Dr. Hamish L. Fraser, Adviser
Dr. James C. Williams
f8~
::e~t
of Materials
Science and Engineering
ABSTRACT
Recent advances in laser deposition technology have made the production o f
advanced materials more technically feasible. By utilizing the unique characteristics o f
the laser deposition process, materials can be made that are difficult to produce by con­
ventional methods. Dispersion hardened titanium alloys hold promise for elevated
temperature structural materials. Production o f these alloy using rapid solidification
processing (RSP) techniques requires an additional consolidation operation that can
destroy the benefits gained though RSP. Laser direct deposition can be used to deposit
these materials near net shape and produce a bulk, dispersion hardened alloy that has the
desired structure for improving properties without the need for further processing. Depo­
sition o f titanium powder with additions o f boron and erbium was accomplished by using
an Optomec LENS™ (Laser Engineered Net Shaping) machine. The resultant micro­
structures showed a homogeneous structure with a fine dispersion o f hard second phase
particles. These deposits were characterized using scanning and transmission electron
microscopy, x-ray diffraction, and by energy dispersive spectroscopy in effort to predict
their mechanical properties. The results show that these alloys likely have tensile
strengths surpassing that o f conventional titanium alloys and they retain their high
strength even after long term exposure at elevated temperatures.
ii
Dedicated to my parents
lll
ACKNOWLEDGMENTS
I would like to thank my adviser, Hamish Fraser, for his guidance and support of
this work. I also want to thank all the other student in the Fraser research group for the
many intellectual discussion and the valuable insight I gained from these discussions.
I would like to thank my fellow workers on Project Lightspeed in the Laser Direct
Manufacturing facility at Lockheed Martin, Fort Worth, TX. I would especially like to
thank Ken Cluck for his help in the lab with sample deposition and preparation.
I thank Su Meng for her tireless help in preparing TEM samples. I also thank
Henk Colijn of the Campus Electron Optics Facility for his help and guidance with all of
the characterization equipment.
I am grateful for the personal and professional guidance I received from Tom
Lienert, Michael Breslin, and Xiao-Dong Zhang.
I thank Amy Shapaker for meticulously proofreading this thesis and making it a
legible, coheret document.
This research was supported by a grant from Lockheed Martin Aeronautics Company, Fort Worth, TX.
IV
VITA
November 10, 1975 ...................................... Born - Centerville, IA
1998 .............................................................. B.S. Metallurgical Engineering, University
of Missouri-Rolla
1998-2000 ..................................................... Graduate Research Associate, The Ohio
State University
Summer 1999, Winter 2000 ......................... Engineering Intern, Lockheed Martin
Aeronautics Company, Fort Worth, TX
FIELDS OF STUDY
Major Field: Materials Science and Engineering
v
TABLE OF CONTENTS
Dedication .......................................................................................................................... iii
Acknowledgments ............................................................................................................. iv
Vita ...................................................................................................................................... v
List of Tables ................................................................................................................... viii
List of Figures .................................................................................................................... iv
Chapters:
1.
Introduction ........................................................................................................... 1
1.1
1.2
1.3
2.
Background.................................................................................................... 1
The Laser Engineered Net Shaping (LENS) Process .................................... 3
1.2.1 Laser System ..................................................................................... 3
1.2.2 Poweder Delivery System ................................................................. 5
1.2.3 Processing Environment .................................................................... 8
1.2.4 Motion Control System ................................................................... 10
Advantages to LENS Processing ................................................................. 10
Strengthening Mechanisms .................................................................................. 13
2.1
2.2
2.3
2.4
Solid Solution Strengthening ....................................................................... 14
Hall-Petch Grain Size Strengthening ........................................................... 20
Orowan Dispersion Strengthening ............................................................... 24
Summary ...................................................................................................... 29
3.
Prior Experimental Work ...................................................................................... 32
4.
Experimental Procedure ....................................................................................... 44
5.
Rapid Solidification .............................................................................................. 49
Vl
6.
Laser Deposited Titanium Alloys ......................................................................... 58
6.1
6.2
6.3
6.4
Commercially Pure Titanium LENS Deposits ............................................
CP-Ti Plus Elemental Boron LENS Deposits .............................................
Prealloyed Ti-8Al-1Er Powder ....................................................................
6.3.l Single Pass Ti-8Al-1Er Beads .........................................................
6.3.2 Ti-8Al-1Er LENS Deposits .............................................................
Conclusions .................................................................................................
Bibliography
58
63
71
74
76
99
....................................................................................................... 101
vu
LIST OF TABLES
2.1
Size mismatch and approxiamte solubility of certain alloying
elements in a-titanium .......................................................................................... 15
4.1
Process parameters used for the various titanium alloy deposits ......................... 45
4.2
Heat treatment schedule for the various as-deposited samples ............................ 47
6.1
Vickers hardness values for the various laser deposited samples ......................... 62
6.2
Chemical analysis of bulk Ti-8Al-1Er deposits .................................................... 93
6.3
Table of experimental and calculated values used in strength
calculations for Ti-8Al-1Er deposits .................................................................... 95
Vlll
LIST OF FIGURES
Figure
Page
1.1
Photograph of Optomec LENS system at Ohio State ............................................. 4
1.2
Normal spectral reflectance versus wavelength for various materials ................... 6
1.3
Schematic of nozzle/laser beam arrangement in the LENS process ...................... 7
1.4
Schematic of the atmosphere/carrier gas flow in the LENS system ....................... 9
2.1
Schematic of hexagonal close-packed lattice structure of a-titanium
showing interstitial sites ....................................................................................... 15
2.2
Binary phase diagrams for Ti-Al, Ti-C, Ti-N, and Ti-0 ....................................... 16
2.3
Mechanical properties of titanium with 0-8 weight percent addions
of aluminum .......................................................................................................... 17
2.4
Influence of various interstitial alloying elements on the shear
strength of titanium at different temperatures ...................................................... 21
2.5
Effect of grain size on strength in titanium alloys with various amounts
of interstitial oxygen ............................................................................................. 22
2.6
Schematic of Orowan bypassing mechanism ....................................................... 25
2. 7
Schematic of dislocation bowing between two hard barriers under
an applied shear stress .......................................................................................... 25
3.1
Binary phase diagrams for Ti-B and Ti-Er ........................................................... 35
3 .2
Strengthening contribution from various hardening mechanism
in Ti-lEr alloy ....................................................................................................... 41
3.3
Strength increment as a function of temperature for Ti-lEr and
Ti-1.5Nd alloys ..................................................................................................... 42
lX
5 .1
Schematic showing isotherm in thin plate with line heat source showing
the adaptation made to account for laser processing conditions .......................... 50
5.2
Graph showing the cooling rate versus thickness based on the modified
equations for heat transfer in welding processes .................................................. 53
5.3
Scanning electron micrograph of bulk laser deposited 316 stainless
steel sample .......................................................................................................... 55
5.4
Scanning electron micrograph of thin wall laser deposited 316
stainless steel sample ............................................................................................ 56
6.1
SE SEM micrograph of as-deposited CP-Ti sample ............................................. 59
6.2
Brightfield TEM micrographs showing grain boundaries in as-deposited
and heat treated CP-Ti samples ............................................................................ 60
6.3
BSE SEM micrograph of as-deposited CP-Ti + lB sample ................................. 64
6.4
X-ray diffraction scans for the CP-Ti and CP-Ti + lB sample ............................ 65
6.5
Brightfield TEM micrograph ofTiB 2 particle and diffraction pattern used
for identification ................................................................................................... 66
6.6
BSE SEM micrograph ofCP-Ti + lB deposit showing microstructure
near a large void ................................................................................................... 68
6. 7
BF TEM micro graphs of the heat treated CP-Ti + 1B sample showing
showing precipitate morphology .......................................................................... 69
6.8
BSE SEM micrograph showing microstructure of as-received Ti-8Al-1Er
powder .................................................................................................................. 72
6.9
BSE SEM micrographs of a single bead deposit ofTi-8Al-1Er ........................... 75
6.10
BSE SEM micrographs of the as-deposited Ti-8Al-1Er structure ........................ 77
6.11
BF TEM micrograph of the bulk as-deposited Ti-8Al-1Er structure ................... 78
6.12
BF TEM micrograph of thin wall Ti-8Al-1Er deposit showing a bimodal
distribution of second phase particles ................................................................... 79
x
6.13
BSE SEM micrograph showing bands of precipitates coinciding with
bands of porosity .................................................................................................. 80
6.14
BSE SEM micrographs of the Ti-8Al-1Er bulk and thin wall deposits
heat treated at 560°C for 30 minutes .................................................................... 82
6.15
BSE SEM micrographs of the Ti-8Al-1Er bulk and thin wall deposits
heat treated at 700°C for two hours ...................................................................... 83
6.16
BSE SEM micrographs of the Ti-8Al-1 Er bulk and thin wall deposits
heat treated at 700°C for 100 hours ...................................................................... 84
6.17
BSE SEM micrographs of the Ti-8Al-1 Er bulk deposits comparing the
size of precipitates in the as-deposited and 100 heat treated condition ................ 85
6.18
BSE SEM micrographs comparing particle size in the thin walled
as-deposited Ti-8Al-1Er and the 100 hour/700°C heat treated sample ................ 86
6.19
Brightfield TEM micrographs showing how the precipitate particles
line up along dislocations ..................................................................................... 88
6.20
Brightfield TEM micrographs comparing the precipitate size and
distribution in the 100 hour heat treated bulk and thin wall
Ti-8Al-1Er deposits .............................................................................................. 89
6.21
Brightfield TEM micrograph of isolated precipitate particle with
simulated and experimental [112] zone axis diffration pattern ............................ 91
6.22
Brightfield TEM micrograph showing preferred orientation particle
coarsening in the 100 hour heat treated Ti-8Al-1Er deposit.. ............................... 92
6.23
Brightfield TEM micrograph showing dislocation/particle interaction ............... 97
Xl
CHAPTER!
INTRODUCTION
1.1 Background
Solid freeform fabrication (SFF) techniques have existed for many years and have
taken on many different forms. Early work using photo-curable plastic resins
(stereolithography) led to the development of similar processes using ceramics and
metals. Recently, direct fabrication of metallic components using SFF has developed into
a viable manufacturing technology [ 1]. This technology has opened the door for the
production of advanced materials that are unique to the process. By taking advantage of
rapid solidification and inert atmosphere control, alloys can be created that have characteristics vastly different from conventional materials. Alloying can be done in-situ using
elemental powder blends, to easily create custom alloy compositions and offer the possibility for creating functionally graded materials. Another key feature of SFF is that the
resultant object is produced near-net shape. This allows for the production of alloys that
do not need further thermomechanical treatment. It is also an important characteristic
when considering brittle and other difficult to process materials. These unique features
can all be achieved by using solid freeform fabrication to create structural metallic materials.
Stereolithography was the first milestone in solid freeform fabrication technology.
In this process, a solid model is created using computer aided design (CAD) software and
is subsequently sliced into very thin planar sections. This information is then used to
1
control a laser as it traces out each slice. A mesh base plate, which is used to support the
part as it is being processed, is placed in a bath of photo-curable plastic resin. A lowpowered ultraviolet laser is then used to selectively cure the liquid resin layer by layer.
After each layer is cured, the mesh plate moves down slightly so that fresh liquid resin is
exposed to the top surface of the previous layer. The next layer is then cured directly on
top of the previous layer. This process is repeated until the entire part has been fabricated. The result is a fully dense, three-dimensional object that is an exact replica of the
computer generated solid model.
An obvious benefit of this process is that new designs can be rapidly prototyped
and easily changed before the part is put into service. A drawback, however, is that the
part is made from a plastic resin that is not suitable for end use applications. Consequently, a mold must be made from the cured resin part and the final object must be cast
from some other structural material. Naturally, this drawback led to the development of
systems that could perform the same process with applicable materials (i.e. metal, ceramic).
Laser cladding was first patented by the United Technologies Corporation in the
late 1970s [2]. This was the first commercial attempt to use a concentrated heat source to
accurately create a solid metallic deposit. This type of process, however, merely created
a layered deposit and was not capable of producing complex structures. Laser cladding
was abandoned as a practical method of achieving a controlled three-dimensional buildup
of material.
The combination of stereolithography and laser cladding brought about the development of a process capable of fusing metal feedstock with a laser in a precise, controlled
manner. The process was developed concurrently at Sandia and Los Alamos National
Laboratories in the early 1990s. Both processes are similar in that they use a focused
laser beam as a heat source to melt metallic powder to create a solid, three-dimensional
2
object. Sandia called the process LENS™ (Laser Engineered Net Shaping) and licensed
it to a private company, Optomec Design Company of Albuquerque, New Mexico.
Optomec sold their first unit to The Ohio State University (Figure 1.1) in April of 1998
and now has systems located throughout the country. The remaining portion of this
chapter will discuss the nature of the LENS process and list some of its key advantages
over conventional processing.
1.2 The Laser Engineered Net Shaping Process
The LENS process can be divided into four main areas: the laser system, the
powder delivery system, the processing environment, and the motion control system.
Each of these areas will be discussed along with major advantages and drawbacks.
1.2. l Laser System
The heat source for the LENS process is typically a solid state, however, systems
exist that utilize higher power gaseous lasers. Depending on the type and power of the
laser, the beam can be delivered via fixed or fiber optics. The most common choice is the
Nd:YAG laser which produces near-infrared radiation at a wavelength of 1.064µm at a
nominal output power of about 700 watts. Typical working power, accounting for losses
through the optics, is between 150 and 500 watts. This energy is focused to a spot of
approximately 0.75mm in diameter, thus giving power densities between 30,000 and
100,000 W/cm 2• As mentioned before, higher power C02 lasers have been used to increase deposition rates. These lasers operate at 10.6µm at powers up to 15,000 watts.
While these lasers can increase the deposition rate by increasing the size of the molten
pool, there are drawbacks that can limit their use. First, high power C0 2 lasers cannot be
delivered with fiber optics. This limits the motion of the workhead to strictly up and
down z-axis motion, whereas a heat source delivered via fiber optics can be manipulated
3
Figure 1.1: Photograph of the Optomec LENS system. The powder delivery system is
on the upper left of the glove box, the laser head is contained in the black compartment
on top of the box, and the laser power supply is to the right.
4
with a robot arm. Systems utilizing a C02 laser are usually limited to 2-112 dimensions
(limited unsupported overhang capability) whereas fiber optic delivered systems can
achieve full three-dimensional capability. Another drawback to the C0 2 laser is the
higher reflectance of common metals at longer wavelengths as seen in Figure 1.2. Although the higher power makes up for the increased loss due to reflectance, the overall
efficiency of the system is reduced.
1.2.2 Powder Delivery System
Metal powder is delivered to the workhead by a motor-driven pneumatic feeder
system. Commercially available screw-driven mechanisms are commonly used for this
process. In the Optomec LENS system, a proprietary powder delivery system is used that
involves a rotating disk instead of a screw. The operating principle, however, remains the
same. The powder is delivered into a moving gas stream (typical carrier gas is argon)
that carries the powder to the workhead. Control of the powder flow can be achieved by
varying the speed of the powder moving mechanism. Powder exit velocity can be controlled by varying the carrier gas pressure.
Once the gas/powder mixture has been delivered into the working chamber, a
splitting device is used to separate the powder into multiple streams. As shown in the
nozzle arrangement in Figure 1.3, four streams are typically used. The four nozzle/90°
arrangement allows for a sound deposit regardless of workhead travel direction. The
laser beam travels down the middle of the nozzle arrangement and is brought to a common focus point with the powder. This common focus is lowered to a metal substrate
upon which the part is deposited. The laser beam creates a molten pool into which the
powdered metal is injected, forming a small metallic deposit. As the workhead is moved,
a line of deposited metal is formed on the substrate. Adjacent lines are laid down to form
a thin layer of a designed part. The z-axis is then incrementally moved upwards and the
5
-- -- --..
-- --..· . ---
1.0
I
aluminum
;
0.8
---- -- - ' ...... ----.... .....
......
Q)
..
.
.
.
.
...
g
.$
(.)
Q)
~
~
(.)
Q)
0.
C/)
0.4
...
....
/
... -'* ..
l/titanium
nickel/
v
~
... ..........
~-·
~·
..-- ..-
_....··"";tee 1- -
/
... ---::: ~
.
.
..
... .. / . v··
.
.. ., /
0.6
~
]
....
...
;:· .....
.
.
.
z0
0.2
0.1
1.0
10 .
Wavelength (microns)
Figure 1.2: Graph showing normal spectral reflectance versus wavelength for various
materials. The LENS process uses a Nd:YAG laser with a wavelength of 1.064 µm.
Other systems use a C0 2 laser at the expense of higher reflectance. Adapted from
Touloukian and Ho [3].
6
(a)
powder
delivery
nozzles
(b)
powder
delivery
nozzles
laser
beam~
Figure 1.3: Schematic showing the nozzle/laser beam arrangement in the LENS process
(a) and a cross-section of the deposition process (b).
7
next layer is deposited on top of the first. This procedure is repeated until the part is
complete.
Measurement and control of the powder feedrate is difficult to achieve. In the
LENS system, three separate variables can affect the powder mass flow rate: the motor
speed of the powder feeder, the back pressure in the powder hopper, and the mass flow
rate of the carrier gas. Variations in any one of these can cause an unknown change in the
flow rate of powder. Attempts have been made to precisely measure the powder flow rate
using a Coriolis meter. Unfortunately, the abrasiveness of the powder destroys the internal mechanical parts of the meter. Currently, the only way to accurately determine
flowrates is to capture powder in a container at various settings and weigh the mass per
sampling time. This empirical method, however, is not very accurate and can vary
considerably on a day-to-day basis.
1.2.3 Processing Environment
The processing atmosphere is typically composed of argon which allows for the
deposition of highly reactive metals without the fear of heavy oxidation. A purification
system is usually employed to help reduce the oxygen content to below five parts-permillion. This level of atmosphere purity is required for deposition of metals that can
dissolve or react with oxygen forming unfavorable phases and/or properties. Titanium,
for example, can dissolve large amounts of oxygen and nitrogen, which can cause brittleness in the material and result in cracking during deposition. Controlling the content of
the atmosphere gas is crucial in forming a sound deposit that has the desired properties.
Argon from inside the chamber is pumped into a recirculation tank. This compressed
atmosphere gas is used to carry the powder into the workhead. Figure 1.4 shows a
schematic of the atmosphere control system. This closed loop system allows for the
addition of metered amounts of alloying gases (i.e. oxygen, nitrogen) into the chamber.
8
atmosphere
gas
to powder feeder
assembly
recirculation
baffle tank
process mg
chamber
flow
meter
shielding gas
powder
feeder
Figure 1.4: Schematic of the atmosphere/carrier gas flow in the LENS system.
9
flow
meter
In addition to the powder delivery gas stream, a stream of shielding gas is introduced to
the workhead to prevent splatter of molten debris onto the focusing lens.
1.2.4 Motion Control System
The first step in making a laser deposited part is to create a solid three-dimensional model of the object using standard CAD software. This computer rendered model
is then sliced into a discrete number of layers using a slicing algorithm. Three-axis
LENS systems typically utilize stereolithography slicing software for the slicing routine.
This type of slicing limits the part to 2-1 /2 dimensions since all of the slices must remain
in the x-y plane. More complex slicing software is needed to create the path plan for
systems capable of more than three axes of motion (i.e. full three-dimensional capability).
Work is currently underway that will optimize the laser head path plan for complex threedimensional structures. Slicing information includes the number of layers, the spacing
between each layer, the spacing between the rows in individual layers, laser head travel
speed, and other pertinent processing information. The interaction among these processing variables is not yet fully understood. Titanium is typically run at about 250 to 300
watts at a travel speed of 12 to 20 inches per minute. These two variables are highly
interdependent, as is the layer spacing and spacing between individual rows. It is critical
to understand these interactions to create a laser path plan for complicated structures.
Although these interactions will likely be verified over time, the optimization of the
deposited structure is, for now, an art rather than a science.
1.3 Advantages to LENS Processing
Laser Engineered Net Shaping technology allows for some attractive benefits in
terms of efficiency and simplicity. First, no molds or dies are needed to create a part,
eliminating many intermediate steps in the manufacturing process. Furthermore, the part
10
is produced near-net shape, no forging or other thermomechanical post treatments are
needed. These benefits reduce the need to procure expensive molds, dies, and forging
blanks and can therefore save considerable time and money. By utilizing LDM, a prototype part can be made quickly, thus reducing the concept-to-product time required.
Unlike machining of large billets, very little product is wasted, further reducing the cost.
Although the efficiency of powder usage is only around 10 to 20 percent, any powder that
is not consumed in the molten pool can be easily recycled and used again.
By exploiting some of the unique characteristics of the laser deposition process,
new materials can be developed that would be difficult to produce using conventional
methods. Rapid solidification can be achieved in LENS deposits. Solidification rates in
the range of 103 to 105 Kl second allow for the formation of fine grained microstructures
with improved mechanical properties. This rapid solidification also allows the formation
of unique metastable structures in a bulk material that could not be attained through
conventional methods. Stable structures are obtained upon heat treatment resulting in
materials with unconventional microstructures and properties.
The feed material for the LENS process in typically powdered metal, although
wire feed has been used successfully. The use of powdered metal allows for greater
flexibility in alloy composition. By blending elemental powders prior to deposition,
alloys with unique compositions can be formed. These elemental blends have been
shown to react in the molten pool and create an in-situ homogeneous alloy. This type of
alloy tailoring can be used to create variable composition alloys or functionally graded
materials. Composite materials can be produced in this manner as well. Work has been
done that blends a ceramic silicide material into a ductile metallic phase thus producing a
remarkably strong high temperature composite [4].
The laser direct manufacturing process, although relatively new, holds tremendous promise for developing exciting new materials. While traditional alloys were
11
developed to fit a particular process such as casting or forging, the unique aspects of the
LENS process, such as rapid solidification and atmosphere control, can be taken advantage to develop new materials with beneficial properties. These process characteristics
can produce an alloy that is strengthened by multiple means. The next chapter will
explore the possible strengthening mechanisms available to LENS processed materials
and estimate their significance based on established equations and prior work.
REFERENCES
1 Keicher, D. M. and W. D. Miller, Metal Powder Report 53 12 (1998).
2
Brown, C. 0., E. M. Breinan, and B. H. Kear, United States Patent# 4,323, 756.
3
Touloukian, Y. S. and C. Y. Ho, eds. Thermal Radiative Properties; Metallic Elements
and Alloys. !FI/Plenum, New York ( 1970).
4
Brice, C. A., K. I. Schwendner, S. Amancherla, H. L. Fraser, and X. D. Zhang, to be
published research (2000).
12
CHAPTER2
STRENGTHENING MECHANISMS
The unique aspects of the Laser Engineered Net Shaping (LENS) process allow
for the creation of new alloys that can be strengthened by multiple mechanisms. The two
dominant process characteristics, rapid solidification and atmosphere control, can be used
to introduce strengthening features into the microstructure. Rapid solidification increases
the solubility of certain alloying elements. These elements can then be precipitated out,
either during the process or by heat treatment, giving a fine distribution of second phase
particles that are ideal barriers to dislocation motion. Rapid solidification can also be
utilized, in conjunction with a grain pinning element such as boron or erbium, to create a
fine grained microstructure that increases the material strength via the Hall-Petch relationship. Atmosphere control allows the precise injection of gaseous alloying elements
such as oxygen and nitrogen into the microstructure. These elements are excellent
strengtheners in controlled amounts and the nature of the LENS process allows the right
amount to be added without the fear of brittleness caused by excessive additions. This
chapter discusses each strengthening mechanism, determines which variables are important in optimizing strength, and applies the results to the LENS process to create novel
titanium alloys.
13
2.1 Solid Solution Strengthening
Solid solution strengthening occurs when alloy atoms are randomly distributed
throughout the matrix, causing local misfit strains that affect dislocation behavior. Solute
atoms can also cluster around dislocation cores, creating "solute atmospheres" that
impede dislocation motion. Elements can occupy either regular sites in the lattice (substitutional) or sites in the interstices (interstitial). Two different interstitial sites exist in the
close-packed hexagonal unit cell as indicated in Figure 2.1. Typically, the type of site
that a particular atom occupies is determined by its atomic radius and electronic character. In a-titanium, most interstitial elements occupy the octahedral sites since their
atomic radii closely match the radii of the octahedral sites (the tetrahedral sites have a
radius that is approximately half that of the octahedral sites). Other alloying elements
with much larger radii (aluminum, tin, etc.) occupy substitutional sites in the lattice.
Table 2.1 shows common titanium alloying elements with their size mismatch and solubility limits in a-titanium. Figure 2.2 shows partial binary phase diagrams for some of
the alloying elements that can contribute solid solution strengthening effects in titanium.
For substitutional alloying, an element with high solubility is desired so that a
large amount of solute can be added without fear of precipitating a second phase. The
optimum choice for titanium is aluminum, which can remain in solution at concentrations
up to 25 wt% (about 8 wt% at 298 K). Figure 2.3 shows the strengthening effect of
aluminum when added to titanium. Addition of aluminum to the base titanium metal
must occur prior to the laser processing operation and hence cannot be controlled during
deposition. On the other hand, gaseous interstitial elements can be added in-situ giving
greater flexibility in the level of hardening obtained. Since control of solid solution
hardening in LENS processed materials occurs through metered gaseous additions of
interstitial elements, the remaining review will focus on interstitial, not substitutional,
solid solution strengthening.
14
(0011
(100)
Figure 2.1: Schematic of hexagonal close-packed lattice structure of a-titanium showing
interstitial sites. Dark circles are tetrahedral sites while small open circles are octahedral
sites. Adapted from Conrad [ 1].
Solute
Approximate
atom diameter
size difference, Approximate
% (d11 - dM)/ solubility in
a-Ti, wt%
dm%
Hydrogen* .....
Carbon* .......
Nitrogen* .....
Oxygen* ......
Magnesium ....
Aluminum* ....
Silicon ........
Vanadium* ....
Chromium .....
Manganese ....
Iron ..........
68
49
51
59
-9
2
19
10
14
23
15
Solute
0.2
0.5
4
12
0.1
25
2
4
<l
1
1
Approximate
atom diameter
size difference, Approximate
% (d11 - dM)/ solubility in
a-Ti, wt%
dTh %
Cobalt ........
Nickel ........
Copper ........
Gallium .......
Zirconium . . . . .
Niobium ......
Tin* ..........
Molybdenum* ..
Tantalum ......
Tungsten ......
14
15
12
16
- 8
2
-3
7
2
6
1
1
1
10
100
4
22
1
l
1
Note. Solutes marked with an asterisk (*) are common additions to Ti.
Table 2.1: Table showing size mismatch and approximate solubility of certain alloying
elements in a-titanium. Adapted from Brooks [2].
15
(a)
0
1800
JO
0
(b)
20
3000
1700
1600
1500
2500
1400
(PTi)
1300
2000
1200
1100
1.8
1000
1500
900
'
(PTi}
I
I
800
I
I
I
I
I
I
I
700
1000
'
'
'
I
600
I
(aTi)
I
500
0
20
JO
30
10
(c)
(d)
2
0
6
4
20
0
2200
3500
2000
3000
L
1800
2500
4.0 , ...
2000
.........
-----
L
-----f5,2
1400
,12.5
,,'',,'
,,,, , /
I/
,
,,
I
1500
,
,,
,,
,
,,
1000
882"C
,,, .
500
0
, ...
5
1600
/
I
I
I
I
I
1200
(aTi}
1000
(aTi}
600
--- -----
_... ...
------
600
400
10
0
15
10
20
Figure 2.2: Parital Binary phase diagrams for (a) Ti-Al, (b) Ti-C, (c) Ti-N, (d) Ti-0.
Adapted from Okomoto et al [3].
16
....
L
~
0
o.~
~.~
tit
10
LJ
0
1
2
3
4
Aluminium
5
6
7
8
0
wt. /o
Figure 2.3: Graph showing mechanical properties of titanium with 0 to 8 wt% addtions
of aluminum. Adapted from Jaffee [4].
17
Conrad [ 1] classified the dislocation/interstitial solute interaction energy into three
main categories. He considered the elastic, chemical, and electrical effects individually
and then calculated a relationship between shear strength and concentration based on the
combination of these three effects.
The elastic effects can be further broken down into size misfit and modulus
mismatch. Most interstitial elements are slightly larger in radius than the octahedral sites
in the hexagonal close-packed titanium structure. Occupation of these octahedral positions in the titanium lattice creates a tetragonal distortion and increases the c/a ratio. This
size misfit affects both the hydrostatic and shear components of a dislocation's stress field
and thus can have a significant interaction with both edge (hydrostatic and shear stress
components) and screw (shear stress component only) dislocations. Modulus mismatch
can have a repulsive or an attractive effect on a dislocation. This effect is similar to the
Orowan hard barrier mechanism (Section 2.3) in that the difference in the elastic modulus
of a second phase influences the energy of a dislocation. Dislocation energy is directly
proportional to the shear modulus of the material it is passing through. A higher modulus
of the second phase means the dislocation will have a more difficult time moving through
it, thus creating a repulsive effect. Similarly, a lower second phase modulus has an
attractive effect by making dislocation travel easier. These modulus mismatch differences contribute to the overall solid solution strengthening effect.
Chemical interactions can also play a part in how individual solute atoms interact
with dislocations. These interactions can be classified into four categories: clustering,
stacking faults, dislocation core, and chemical bonding. Clustering of solute atoms
increases the size and decreases the distribution resulting in a reduction in dislocation
interaction. However, at the low interstitial contents considered in this work ( < 1 at%),
clustering is not expected to occur [ 1]. Stacking faults can also have a chemical interaction with dislocations. The formation of a stacking fault produces a region of different
18
crystal structure and thus can have an effect on how the solutes behave (i.e. solubility
differences in various crystal structures, radii of interstitial sites, etc.). Because titanium
has a low stacking fault energy, this type of interaction is not expected to have much
effect. The core structure of the dislocation in a close-packed hexagonal material has
been shown to have a different crystal structure than that of the matrix [1]. This effect
can increase the misfit of the solute atoms in a way similar to a stacking fault. This
increase in energy leads to a repulsive action of the solutes in the dislocation core. Finally, chemical bonding can play an important role in the interaction between solutes and
dislocations. Strong directional bonds may form between the solute and the lattice
material. These high energy bonds require a large stress to break. Therefore, the passing
of a dislocation through this higher bond energy requires more force and can lead to an
overall increase in strength.
Electrical interactions result from changes in the electronic structure of atoms in
solution. Very little work has been done to investigate these interactions. Cottrell et al
[5] have shown that electrons can form dipoles when in the vicinity of an edge dislocation. The valence electrons can behave like a gas and tend to segregate away from the
•
compressive side of an edge dislocation. Thus, the solute atoms distribute themselves in
a favorable position along a dislocation. Though this effect has been shown to occur in
copper alloys, it is insignificant compared to the other strengthening mechanisms imparted by randomly distributed solute atoms.
Conrad has proposed a fairly simple model for the effect of a dilute concentration
of solute atoms on the strength of titanium. His relationship states that the shear strength
contribution from interstitial solute elements,
r0
t ,
= µ:C[1+(sc)1]
0
is
2.1
where µ is the shear modulus, sis the size misfit parameter and C is the concentration in
19
atomic percent. This relationship assumes that all the barriers have the same strength,
which is not always the case. Conrad goes on to show that the addition of weak barriers
is given by
2
Tweak
= T1 2 + T2 2
2.2
and the addition of mixed barriers is given by
T mixed
= Tstrong + Tweak
2.3
For a random distribution of individual solute atoms, it is difficult to predict strong versus
weak behavior. Equation 2.1 will be used throughout this text as an estimation of the
strengthening effect of interstitial solute atoms in the titanium matrix. Comparisons will
also be made with the experimental results shown in Figure 2.4
2.2 Hall-Petch Grain Size Strengthening
Grain size strengthening is a well known method for controlling and improving
the mechanical behavior of alloys. As the grain size of a polycrystalline sample decreases, the strength increases. This typically holds true for grains in the micron size
range but not always for grains in the nanoscale range [6]. The relationship between
grain size and strength was first developed by Hall [7] and Petch [8]. They showed that
2.4
where cry is the yield stress, k is a constant known as the Hall-Petch slope, dis the average
grain diameter, and cr0 is the stress required to move a dislocation through the matrix (i.e.
the frictional stress). This frictional stress is related to the critical resolved shear stress by
2.5
where Mis the Taylor factor. The Taylor factor, found by averaging the Schmid factor
over all possible orientations, is equal to 3.06 for a randomly oriented sample [9]. The
relationship in Equation 2.4 has been experimentally verified many times for materials
where grain size is the dominant strengthening mechanism. Figure 2.5 shows the grain
20
a-Ti
12
... =r 10-•s •1
Polycrystols ( O'j l
Conrod et al
a Finloy a Snyder
A Jaffee et al
+ •
1000
o •
77
Okazaki et ol
{1oio} < 11.20> GlideC2.51"CRssl
BOO
~z
::E
o •
" •
Elssner et al
Tanaka 6 Conrod
6
ae
C\J
g
b- 400
16
(O+N+C)(ot.%)
Figure 2.4: Graphs showing the influence various interstitial alloying elements have on
the shear strength of titanium at different temperatures. Adapted from Jaffee [4].
21
Q -Tl
300 K
i. 5110-•,-·
0 • 0
A
A•
A • 70 Ti (l.O at. % Oeq.l
0
100
-...
•I
Conrod et al (Wire I
J - IS!leet)
OllOIOlli ., a I (Wirw)
100
z
ll
N
8
tf
(et MARZ Z.ll}
(O)Batt...
C0.2 ot. % 0941
1200
1400
Figure 2.5: Graph showing the effect grain size has on the strength of titanium with
various amounts of interstitial oxygen. Note that the concentration of oxygen has no
effect on k, the Hall-Petch slope. Adapted from Conrad [l].
22
size dependent strength for unalloyed titanium. For dispersion hardened materials,
however, this equation may not be valid. Dislocations interacting with the dispersed
second phase provide most of the strengthening in these types of alloys. Special consideration must be made to account for the dislocation-dispersoid interaction that occurs
inside individual grains.
Initially, it would appear that an adjustment in cr0 would compensate for the
increased strength due to precipitation hardening. That is, the stress required to move a
dislocation through the dispersion hardened matrix could take the place of cr0 for an
unmodified crystal and the Hall-Petch relation would remain the same. Lasalmonie and
Strudel [10], however, show that k can also increase with the increasing strength of the
matrix. Thus, the strengthening increase due to the dispersoids cannot be isolated in the
cr0 term, but must also appear in k.
This dependence of the factor k on a dispersion of second phase particles has also
been demonstrated by Mangen and Nembach [ 11 ]. They used the y' precipitates in
NIMONIC PE16 to show the effect a dispersed second phase has on Hall-Petch grain size
strengthening. In their analysis, they utilize the work hardening model for Hall-Petch
where dislocation interactions within each grain govern the strengthening, as opposed to
grain boundary pileups (i.e. looping, not shearing, occurs). From this analysis, they show
that
2.6
where Lis the average slip length. Hansen and Ralph [12] showed that Lis related to the
grain diameter din oxide dispersion hardened copper. From this, Mangen and Nembach
conclude that k will depend on the strength and spacing of the precipitates since Lis
related to d. This now excludes the linear addition of strengthening mechanisms because
grain size strengthening is dependent on the precipitation strengthening mechanism.
23
To further complicate things, work has been done that shows the effect grain size
has in dispersion hardened alloys may be insignificant compared to the overall strength of
the alloy. Kim and Griffith [13) showed that in precipitate/dispersion hardened aluminum
7091, the strength in the underaged condition was grain size dependent only for grains up
to 50µm in size. For the peak and overaged conditions, the yield strength is independent
of grain size for grains in the micron size range.
Dispersion hardening elements, when added to titanium, can also produce a
pronounced grain pinning effect. Thus, the grain size in these alloys produced via rapid
solidification is expected to be small and the Hall-Petch mechanism cannot be ignored.
Grain size strengthening will be considered in this work, but, it will be addressed as being
secondary to the primary mechanism of dispersion hardening.
2.3 Orowan Dispersion Strengthening
Rapid solidification during the laser deposition process allows the formation of
supersaturated solid solutions that can subsequently be heat treated to form a dispersed
second phase. This second phase is homogeneously distributed throughout the matrix
and provides a means for strengthening via the Orowan relationship, provided these
particles are non-shearable. These nanoscaled precipitates interfere with the motion of
dislocations and thus increase the stress required for a given strain. Figure 2.6 shows
schematically how precipitates can cause an increase in strength in a dispersion hardened
alloy. The dislocations are not able to pass through the second phase material but the
section of the dislocation line that is free to move continues to do so. Eventually, the
dislocations pinch off on the other side of the particles and leave behind dislocation
loops. As seen in the diagram, this has the effect of decreasing the mean particle spacing,
A, which can have a direct effect on the Orowan strengthening mechanism. In the present
24
9
©
~
©
A,
~
A,'
©
Figure 2.6: Schematic diagram showing Orowan bypassing mechanism. The dislocation
line is bent around the particles as it moves through the matrix until it is pinched off
leaving Orowan loops around the particles.
a)
tF
Figure 2.7: a) Schematic of dislocation bowing under an applied shear stress in theydirection. The line tension, T, is the equilibrium maintaining force. b) Schematic
showing angle cp used to determine critical Orowan stress.
25
study, the matrix material is titanium and the dispersed second phase is achieved by the
addition of a compound forming alloying element. These particles form hard barriers that
are incoherent with the matrix and provide the ideal Orowan obstacle.
The derivation of the Orowan equation is fairly straightforward. Fig. 2.7a shows
the balance of forces on a dislocation under force F that is pinned at two hard barriers
lying in a common slip plane. The mobile portion of the dislocation tends to bow into a
semicircular shape while the ends remain pinned at the particles. The closer together
these barriers are, the greater the effect of the strengthening mechanism. To prove this
dependence, the Orowan relationship is derived beginning with the generalized PeachKoehler equation for force per unit length (FIL) of a dislocation. In the equation
F = i( G/;3 - G/;2 ) - j( G,(3 - G3( 1) + k( G1( 2 - G2( 1)
2.7
L
and Gk =bp;k where bis the burgers vector, (is the line direction, and k = 1, 2, 3 and for
each k, i = 1, 2, 3. Simplification, after consideration of only the forces that do work (i.e.
in the direction of the burgers vector b ), gives
F = J(G3(1) = }bp·23
L
Let r = shear stress in they direction (i.e. cr2 ) . So
F,; = rb}.,,
2.8
2.9
where}.,, is the projected length of the dislocation (the distance between barriers). Now
consider a dislocation pinned between two hard barriers. The energy per unit length of
dislocation can be equated to the line tension, T, of the dislocation
T= µb2
2
2.10
Summing forces in they direction gives
LP,;= -2T + rb}.,, = 0
2.11
and at equilibrium,
2.12
2T= rbA,
26
2(µ;')=rbA
2.13
µb
r=-
2.14
A
giving the final Orowan stress, r, in terms of the shear modulus, the burgers vector, and
the mean particle spacing. Care must be taken when using this equation because there are
assumptions that can be overlooked, causing error in calculations. The first assumption is
that there is no volume change upon precipitation of the second phase. If a volume
change occurs, there will be elastic strain introduced to the matrix surrounding the particles. This strain field will interfere with approaching dislocations, making the obstacles
appear larger than they really are by not allowing them to actually come in contact with
the particles. The obvious effect of this would be to decrease the effective particle spacing, A, making the actual shear stress higher than the Orowan stress. Also, a difference
in shear modulus between the matrix and the precipitate can have the same effect, provided the difference, L1G = GP - Gm' is positive. This effect is understood by considering
the energy of a dislocation, which is proportional to µb 2• If the shear modulus of the
precipitate is larger than the matrix, there will be a repulsive effect once the dislocation's
strain field comes in contact with the particle. When the difference is negative, the
opposite effect can occur and the dislocation can be drawn in towards the precipitates.
For both of these conditions, however, the impact of added stress is minimal. Ashby [ 14]
calculated the shear mismatch standoff distance to be 0.05d, where dis the diameter of
the particle. This factor is too small to be of concern when calculating the Orowan stress.
Equation 2.8 is viewed as being overly simplistic and has been modified numerous times to improve the correlation with experimental data. One of the key problems
with Equation 2.8 is that it does not distinguish between dislocation energies of edge or
screw character. This can be significant since, in general, edge dislocations require less
stress than screw dislocations to move through the lattice. Thus, the Orowan stress will
27
be dependent on the character of the dislocation. Ashby showed that the dislocation
character dependent Orowan stress could be represented by
r=A(e) µb lnd
2n).,
2.15
r0
where rn is the inner cutoff distance (typically taken as 4b) and A(8)
= 1 or 1/(1-v) for
edge and screw orientations respectively. The logarithmic energy term shows the dependence of the stress on the inner (rn) and outer (d) cutoff distances of the dislocation.
Ashby modified this term, using d as the outer cutoff radius instead of A, based on dipole
formation of dislocations around the barriers. Figure 2. 7b shows how the angle <pis
calculated between two bowed-out dislocation sections. The critical angle for hard
barriers is achieved when <p = 0°. This arrangement results in the formation of a dislocation dipole which now allows for interaction among the bowed out dislocations. Ashby
goes on to show that the critical angle is not <p = 0° but occurs somewhere between 0°
and 30°. Computer simulations by Foreman and Makin [15], conducted to analyze the
interaction between bowed out dislocations, confirms this argument.
Consideration of the random distribution of particles in the matrix leads to a
statistical factor, calculated by Foreman and Makin. Assuming that <p = 0°, they calculated a statistical correction factor based on computer simulations. This statistical factor
is due to the variance of the particle spacings in the matrix. Some spacings allow dislocations to easily pass while others do not. In a random distribution, this effect causes a
reduction in the critical Orowan stress. The result is
2.16
cr =Sr
where Sis the statistical correction factor. Foreman and Makin calculated the value of S
to be 0.81 while Kocks [16] calculated it to be 0.84, both in fairly good agreement with
each other. As mentioned previously, the critical breaking angle is not zero, as assumed
in the computer simulations. However, when an increase is made in <p, the statistical
28
factor, S, decreases. These effects trade off with each other resulting in a best approximation for the Orowan stress to be that defined in Equation 2.9. Hirsch and Humphreys [17]
calculated a single Orowan term that incorporates the geometric mean of the two variations of Equation 2.9 and also includes the statistical correction factor. The result
r
= 0.81
(d)
µb
1 ln 2n/l.,(l-v)2
r0
2.17
will be used throughout this paper.
The mean particle spacing, A, must be determined experimentally. The difficulty
in performing this measurement arises from the fact that the best method, transmission
electron microscopy (TEM), is not entirely accurate. An image taken using the TEM
shows a two-dimensional projection through a three-dimensional object. Thus, measuring particle spacings on a TEM negative can result in geometric errors. To account for
this, an approximation is used, again based on the computer simulations of Foreman and
Makin. They have determined that the mean particle spacing is
2.18
where N is the number of particles per unit area of slip plane. The factor 0.81 has already
been included in equation 2.11 so it is ignored in the calculation of A. It is assumed that
any given plane through the sample will give an equivalent average number of particles.
Based on this assumption, any given section analyzed in the scanning electron microscope (SEM) will provide a legitimate approximation of the number of particles on any
given slip plane.
2.4 Summary
Laser direct manufactured titanium can be strengthened in four distinct ways.
First, the addition of a substitutional solute to the base metal results in solid solution
strengthening. This method of hardening is not controllable in-situ, therefore, a given
29
quantity must be added to the alloy prior to deposition. The predicted effect of this alloy
addition can be estimated based on previous work and will be considered a constant
effect in all deposits containing these additions. Variable effects can be introduced by the
addition of gaseous elements through atmosphere control during deposition. These
elements segregate to the interstitial sites and also strengthen the alloy via solid solution
strengthening. Control of these additions, however, can be done in-situ and thus the
mechanical properties can, to some degree, be controlled. The effect of these additions
will be estimated using Equation 2.1 and will also be correlated with previously collected
data. Grain size control is achieved through rapid solidification and additions of grain
pinning elements. These strengthening effects will also be investigated in terms of
known relationships (Equation 2.4) and previous experimental data. The focus of this
paper, however, will be on the effects alloy additions have in creating a fine dispersion of
hard barriers. This mechanism will likely have the largest impact on the overall strength
of the alloy. The experimental results will be evaluated by utilizing Equation 2.17, the
modified Orowan relationship. Finally, all of these strengthening mechanisms will be
compiled and compared with laser deposited alloys of various compositions to develop a
means for roughly estimating the strength of these alloys based on fundamental theoretical considerations.
REFERENCES:
1. Conrad, H., Progress in Materials Science 26 (2-4) (1981).
2. Brooks, C. R., Heat Treatement, Structure and Properties ofNonferrous Alloys.
American Society for Metals. Materials Park, OH (1982).
3. Okamoto, H., P. R. Subrananian, and L. Kacprzak, eds. Binary Alloy Phase Dia
grams, Second Edition. ASM International, Materials Park, OH (1990).
30
4. Jaffee, R. I., in Progress in Metal Physics, B. Chalmers and R. King (eds), Pergamon
Press, New York (1958).
5. Cottrell A.H. and F. R. N. Nabarro, Phil. Mag. 44 (1953).
6. Suryanarayana, C., D. Mukhopadhyay, S. N. Patankar, and F. H. Froes, J Mater Res.
7 (8) (1992).
7. Hall, E. 0., Proc. Roy. Soc. B, 64 (1951).
8. Petch, N. J., J Iron Steel Inst. 174 (1953).
9. Van Aken, D. C.,private communication, University of Missouri-Rolla (1997).
10. Lasalmonie, A. and J. L. Strudel, Journal ofMaterials Science 21 (1986).
11. Mangen, W. and E. Nembach, Acta Metal/. 37 (5) (1989).
12. Hansen, N. and B. Ralph, Acta Metal/. 34 (1986).
13. Kim, Y. W. and W. M. Griffith, Metal Powder Report 1 (1985).
14. Ashby, M. F., in Physics of Strength and Plasticity, A. S. Argon ed. MIT Press,
Cambridge, MA 1969.
15. Foreman, A. J.E., and M. J. Makin, Phil. Mag. 14 (1966).
16. Kocks, U. F., in Physics of Strength and Plasticity, A. S. Argon ed. MIT Press,
Cambridge, MA 1969.
17. Hirsch, P. B. and F. J. Humphreys, in Physics of Strength and Plasticity, A. S. Argon
ed. MIT Press, Cambridge, MA 1969.
31
CHAPTER3
PRIOR EXPERIMENTAL WORK
Producing novel titanium alloys through rapid solidification processing (RSP) is
not unique to the laser engineered net shaping process. Fine grained, dispersion hardened
titanium alloys have been produced utilizing splat quenching, ribbon casting, and other
rapid solidification techniques [1]. These processes, however, produce a powder-like
material that requires an additional consolidation operation. This reduces the benefits of
rapid solidification and also inhibits the practical of these materials outside of a laboratory environment. On the other hand, the laser engineered net shaping process can
achieve rapid solidified structures in a bulk material, without the need for consolidation.
By eliminating these processing steps, a RSP alloy can now be produced both cheaper
and faster, making it more attractive for practical use.
Using RSP to produce conventional titanium alloys does not yield substantial
improvements in mechanical properties. The benefit of RSP in creating high strength
titanium alloys is realized through additions of unique alloying elements;, creating new
alloys altogether. These additions can be dispersoid formers, compound formers, eutectoid formers, or grain pinning agents. This chapter will review some of the previous
work that has been done using RSP to create novel titanium alloys, focusing on dispersion strengthened alloys.
Although the processes used to achieve rapid solidification vary widely, they tend
to give similar results in terms of solidification rates and microstructures. All achieve
32
solidification rates of 103 to 106 K/sec and can produce extremely refined, metastable
microstructures. The benefit of such rapid cooling is an increase in the solubility of most
alloying elements, forming supersaturated solid solutions. These alloys can, upon heat
treatment, produce thermally stable structures that retain high strength levels even at
elevated temperatures. These additions are typically metalloids, rare earth elements, or
strong intermetallic compound formers.
Additions of metalloid elements are known to be effective grain pinning agents
for conventionally processed materials. When added to RSP titanium, these elements
produce a fine dispersion of second phase particles that interfere with dislocation motion
and increase the strength of the alloy. Small additions of the metalloids carbon and boron
tend to produce compounds with the base titanium metal since they have a large electronegativity difference and a high negative heat of solution. Carbon, when in excess of
the extended solubility range, forms a large volume fraction of spherical dispersoids of
the form TiC. Boron also produces a compound with titanium, TiB, but the morphology
of the dispersoids is needle or rod shaped. Figure 3.1 shows a section of the Ti-B phase
diagram.
Rare earth elements are known to scavenge interstitial oxygen when added to
titanium. Small amounts have been added to conventionally processed titanium to remove oxygen from solution and increase the ductility of the alloy [2]. Under RSP conditions, however, the rare earth solutes exhibit extended solid solubility and can be precipitated out to form ideal hard barriers. Rare earth elements have the opposite characteristics in terms of electronegativity and heat of solution when alloyed with titanium. Thus,
they tend to produce an oxide phase and typically do not react with the base metal.
Because the large size of the rare earth elements limits their solubility to substitutional
sites, they have a much lower equilibrium solubility in titanium than the metalloids.
Figure 3.1 shows the Ti-Er phase diagram.
33
Whang conducted a fairly thorough investigation into the effects of various
alloying elements on RSP titanium [3]. Depending on the alloying element involved, the
as-quenched microstructures contained a supersaturated solution of solute and/or solute
clusters. For the rare earth elements, these solute atoms segregated to the subgrain
boundaries only when the solute atom was in excess. For lean solute compositions, the
subgrain boundaries were not seen. Whang also notes that the grain size is nearly inversely proportional to the rare earth solute content. This shows how these fairly small
alloying additions can pin grain boundaries and refine the microstructure.
Heat treatment of these supersaturated metastable structures produces various
intermediate structures, depending on the duration and temperature of the heat treatment.
Heat treatment in the 500-700°C range produces structures similar to G.P. zone precipitation with the combined disappearance of martensite and the sub-grain boundaries. This
precipitation behavior is seen in all RSP alloys with dispersion forming additions. There
are differences, however, in the size, distribution, and morphology of the precipitated
phase. Alloys containing boron produce rod-like precipitates while erbium-bearing alloys
produce spherical precipitates. High temperature heat treatments (above 900°C) affects
the metalloid-containing alloys more significantly than the rare earth-bearing alloys.
While Ostwald ripening is apparent in the boron alloys, very little coarsening is seen in
the rare earth alloys. At these high temperatures coarsening can be expected; 900°C is
above the beta transus, thus producing a more open crystal structure (body centered cubic
vs. hexagonal close-packed) that greatly increases the diffusion kinetics of the solute
atoms.
The mechanical properties of these dispersion hardened alloys is significantly
better than those of conventional alloys. There is, however, quite a difference in the
strengthening effect of metalloid additions versus rare earth additions. A small addition
of metalloid (1 at%) increases the hardness in the as-quenched condition by 4-13%,
34
5
0
10
20
3000
2500
u
0
QJ
....
;::I
....,
"'....a.
2000
QJ
E
QJ
E-
1500
(PTi)
TiB
1000
0
(aTi)
500
0
20
10
40
30
60
50
Weight Percent Erbium
0 10 20 30
40
50
1700
70
60
80
90
100
L
1529°
1500
1400
;;:>
QJ
....
1320 ± 20°c
1300
....,;::I
"'....
a.
1200
E
1100
QJ
(Er
(PTi)
QJ
E1000
0
900
(aTi)
800
700
0
Ti
10
20
30
40
50
60
Atomic Percent Erbium
70
80
90
100
Er
Figure 3.1: Partial binary phase diagram of the Ti-B system (top) and phase diagram for
the Ti-Er system (bottom). Adapted from Okamoto et al [4].
35
whereas the same amount of rare earth raises the hardness 20-4 7%. These effects in the
as-quench (supersaturated solid solution) condition are directly related to the atomic size
mismatch of the alloying elements with titanium. It is also noted that increasing the rare
earth content above 1 at% causes a reduction in hardness.
In the aged condition, the mechanical properties of the alloys are greatly increased
due to precipitation of hard dislocation barriers. The hardness of the alloys peaks in the
intermediate temperature range (500-700°C) and decreases when aged above the beta
transus (approximately 900°C). Again, this can be related to the higher degree of coarsening when held in the beta phase region.
Sastry et al looked at three different titanium-rare earth alloys produced by RSP
[ 1]. Conventionally processed material was compared with the same alloy produced via
RSP. The alloys consisted of small amounts of neodymium, erbium, and dysprosium.
The as-quenched flakes were aged at 700-800°C to precipitate out the dispersion hardening phase from solid solution. All three showed that the dispersoids were refined by an
order of magnitude when produced via RSP. Two weight percent addition of erbium to
titanium produced a homogeneous distribution of incoherent precipitates that increased
the yield strength by as much as 50% without any loss in ductility. Sastry also showed
that a combination of hardening effects can be achieved in these alloys by adding conventional alloying elements. The addition of aluminum to the erbium-bearing alloy (Ti-8Al2Er) produced a fine dispersion of non-shearable erbium-containing particles along with
Ti3Al (a 2), a shearable precipitate.
Sastry also examined the effect of a small addition of metalloids on the microstructures ofRSP titanium. One percent additions of boron and carbon produced a fine
distribution of dispersoids that had a tremendous effect on the mechanical properties of
the alloys. RSP titanium with 0.5 wt% boron increased the modulus by 30% and the
yield strength by 100%. The metalloid alloys can also be combined with the rare earth
36
elements to produce an alloy that takes advantage of both grain size and dispersion
strengthening.
It is important to note that in these experiments, a hot consolidation operation was
performed to put the as-quenched powder in a form suitable for mechanical testing.
Sastry et al vacuum hot pressed the powdered material at 825-900°C for 6 to 16 hours.
This type of treatment can destroy some of the beneficial effects achieved through RSP.
Konitzer et al also experimented with titanium alloys containing erbium additions
[5]. Laser surface melting was used to achieve rapid solidification in a titanium alloy
containing 0. 7 at% erbium. The as-quenched structure contained a small volume fraction
of precipitates in a martensitic titanium matrix. After heat treatment at 700°C for ten
hours, the erbium precipitated out forming spherical dispersoids on the order of 15 nm in
diameter. Electron diffraction was used to identify the precipitate phase. The erbium
particles reacted with the interstitial oxygen in the titanium matrix to form Er20 3 • Furthermore, the orientation relationship was found to be {0001} Ti
II {111} Er20 3 and
<llLO>Ti II <110>Er20 3•
Konitzer also examined the solute controlled coarsening rate, considering both for
excess solute (more erbium than can react with the interstitial oxygen) and for deficient
solute. The classical coarsening theory of Lifshitz, Slyozov, and Wagner was used to
predict the thermal stability of the precipitate particles. In the excess erbium case, coarsening is expected to increase rapidly above the beta transus temperature, while in the
excess oxygen case, the structure is predicted to be fairly stable. Experimental results
confirmed that the excess oxygen microstructures were more stable at higher temperatures when compared with the excess erbium microstructures.
Sastry et al have conducted the most comprehensive study of the effects of rare
earth alloying additions to RSP titanium [6, 7]. Samples were prepared via RSP that
contained various amounts of the rare earth elements Ce, Dy, Er, Gd, La, and Y. These
37
splat quenched flakes were then annealed for various temperatures and times to determine
the effect each element had on precipitation formation. Although many different compositions were analyzed, this review will focus on the erbium addition since it is relevant to
the experimental work described in this paper.
Of the systems investigated by Sastry et al, the most promising with respect to
stable precipitate formation was the erbium-containing alloys. The as-quenched erbiumcontaining samples showed a single phase martensitic structure with all the erbium
dissolved in solution. The Ti-Er alloy showed a distribution of fine particles (average
diameter of 139 nm) after a long term anneal (87.5 hours) at 700°C. Of the titanium-rare
earth alloy systems investigated, the Ti-Er alloy had the smallest average particle size and
was the most stable at sustained high temperatures.
Analysis was done using various characterization techniques to determine the
chemistry of the dispersoid particles. Crystallographic analyis via electron diffraction
was conducted on thinned samples. The results, however, were inaccurate due to unavoidable beam interaction with the matrix. Intensity ratios (Ti:Er), calculated for the
erbium-containing alloys, ranged between 1.2: 1 and 6.1: 1, indicating that the particles
may have a variable composition. X-ray diffraction yielded peaks from the larger particles corresponding to rare earth oxide particles of the form RE 20r Additional peaks
were found, however, that did not correspond to titanium, elemental rare earths, or rare
earth oxides. Electron diffraction confirmed that these small dispersoids (<100 nm) were
not any of the known phases of titanium, rare earth, or their oxides. Auger spectra
showed that the particles were enriched with oxygen and carbon relative to the matrix.
Sastry concluded that these smaller precipitates could not be accurately identified and
referred to them as titanium-rare earth oxycarbides. This is in direct disagreement with
Konitzer et al, who unequivocally identified the particles as rare earth oxides of the form
RE 20 3 for the erbium-bearing alloys.
38
Sastry et al also conducted mechanical tests on these RSP titanium-rare earth
alloys. The splat quenched flakes were consolidated by vacuum hot pressing and were
isothermally forged, rolled, and annealed to produce fully dense, 100% recrystallized
microstructures. The results of the mechanical testing showed that the increase in
strength is dependent on the heat treatment. The highest strength increment is found in
the samples annealed for 100 hours at 700°C, corresponding to a fully aged structure.
These samples, however, show a significant reduction in ductility. The optimum properties are obtained when the samples are annealed for two hours at 700°C. For the Ti-lEr
alloy, the two hour heat treatment produced a tensile strength of 501 MPa (35% increase
over commercially pure titanium) and an elongation of 24.1 % (26% reduction vs. CP-Ti).
For the 100 hour anneal the tensile strength increases to 573 MPa and the elongation is
15.6%.
Sastry et al claim there are three main sources at work in strengthening these
alloys: Orowan strengthening, solid solution hardening, and Hall-Petch grain size
strengthening. An effort was then made to determine the contribution from each mechanism to the overall increase in strengthening. For the Orowan contribution, an equation
similar to that given in Equation 2.11 was used to calculate the strengthening contribution
from hard barriers, while Equation 2.12 was used to calculate the mean free path between
the barriers. The calculated Orowan contribution in the Ti-lEr alloy was 205 MPa for the
two hour anneal and 286 MPa for the 100 hour anneal. The other strengthening contributions were considered with respect to some of the effects resulting from the formation of
the precipitates. These precipitates scavenge the oxygen out of solution, thus reducing
the solid solution strengthening effect. The grain size strengthening contribution can be
influenced by the dispersion of particles within the individual grains. Finally, some
erbium could possibly remain in solution and contribute to the solid solution strengthenmg. Empirical data was used to estimate the contribution from solid solution strengthen39
ing of interstitial oxygen. It should be noted however, that the authors ignored the effects
of nitrogen and carbon in solid solution which could add a significant contribution to the
overall strength of the alloy. Grain size strengthening, also estimated from empirical
data, was expected to have a weak effect due to dislocation-dispersion interactions. The
individual strengthening contributions with respect to oxygen concentration are shown in
Figure 3.2. It is interesting to note that the strength increment appears to be temperature
dependent as shown in Figure 3.3, with the peak increment found around 500°C.
Sastry concludes that acceptable agreement is found with the experimental results
when the three contributing effects are superimposed. This conclusion is based on the
assumption that a maximum amount of oxygen is scavenged for the formation of the
dispersion compound. For the Ti-lEr alloy, the two hour anneal had a measured flow
stress of 414 MPa compared to a calculated flow stress of 404 MPa. For the 100 hour
anneal, the measured stress was 487 MPa compared to the calculated stress of 509 MPa.
Sastry goes on to say, however, that the contributions may not be very accurate estimations based on the actual chemistry of the matrix and particles. It is likely that the errors
from each contribution were such that they negated each other, resulting in a calculated
flow strength similar to that seen experimentally.
The prior experimental work shows many similarities as well as some key differences. Through rapid solidification, normally insoluble alloying elements can be put into
extended solid solubility in titanium. Upon heat treatment, these solute elements precipitate out forming a fine dispersion of hard barriers ideal for strengthening via the Orowan
mechanism. Metalloid elements form titanium compounds when precipitated from the
matrix. Rare earth elements scavenge interstitial oxygen from the matrix to form rare
earth oxides. Positive identification has been made of these particles in Ti-Er alloys by
one group of researchers while another group claims to have found an indeterminate
structure. Thermal stability of these microstructures is best in the Ti-Er alloys and is
40
110
(a)
JOO
90
600
80
-
--
70 ·;;;
ell
0.
~
.......
~
60
::
u 400
so
!::
"'
'O
u
40
>=
"'
u
"'
!::
"'
'O
u
>=
30
200
20
I
ol
0
10
0
(a)
800
(b)
110
JOO
90
600
80
-
70 ·;;;
~
~
~
"'
60
400
so
"'
"'~
v;
'O
'ij
40
>
30
200
20
10
0
o.o
0.1
0.2
0.3
0 ..5
0.6°
11
(wt07o oxygen) ~
Figure 3.2: Graphs showing the contribution from each strengthening mechanism in TilEr. crm is the matrix yield strength, crct is the dispersion strengthening contribution, and
crg is the grain size strengthening contribution. (a) shows results from a two hour anneal
at 700°C and (b) shows results for a 100 hour anneal at 700°C. Adapted from Sastry et al
[7].
41
Ti-1.0Er
Ti-1.SNd - -
--
80
~
~
eu
...
60
C.>
c
·--=
co 40
cu
...
til
Strength increment
(O/o)
20
100
200
::; (
Oalloy -
OTj)
x 100
uTi
300
400
500
600
700
Temperature (°C)
Figure 3.3: Graph showing strength increment as a function of temperature. Note the
peak strength increment at 500°C. Adapted from Sastry et al [7].
42
optimized when there is excess oxygen with respect to the erbium. Significant coarsening occurs above the beta transus due to higher diffusion kinetics in the bee structure.
More importantly, the
a/~
interface creates an accelerated diffusion path for the second
phase particles. Progression of this interface results in extreme coarsening in these
dispersion harded alloys.
REFERENCES
1. 1Sastry, S. M. L., T. C. Peng, P. J. Meschter, and J. E. 0 'neal, Journal ofMetals 35
(9) (1983).
2. Rath, B. B., B. A. MacDonald, S. M. L. Sastry, R. J. Lederich, J.E. O'Neal, and C.R.
Whitsett, Proc. of 4th Internation Conj. on Titanium 2 (1980).
3. Whang, S. H., Journal ofMetals 36 (4) (1984).
4. Okamoto, H., P. R. Subrananian, and L. Kacprzak, eds. Binary Alloy Phase Dia
grams, Second Edition. ASM International, Materials Park, OH (1990).
5. Konitzer, D. G., R. Kirchheim, and H. L. Fraser, in Rapidly Solidified Metastable
Materials 28 (1984).
6. Sastry, S. M. L., P. J. Meschter, and J. E. O'Neal, Met. Trans. A 15A (1984).
7. Sastry, S. M. L., T. C. Peng, and L. P. Beckerman, Met. Trans. A 15A (1984).
43
CHAPTER4
EXPERIMENTAL PROCEDURE
Three separate powdered materials were deposited using Laser Engineered
Net Shaping technology. Two powders, commercially pure (CP) titanium and Ti8Al-Er, were obtained from Crucible Research, Pittsburgh, PA. Both powders were
gas atomized and sieved for a certain size fraction. The nominal size range for the
CP titanium was -100/+325 mesh while the Ti-8Al-1Er powder had a nominal size
range of -80/+200 mesh. Elemental boron powder was obtained from Alfa Aesar,
Ward Hill, MA. The boron powder was nominally-60 mesh and was 99.5% pure
crystalline boron. This powder was mechanically milled and therefore had an
irregular, granular shape.
Deposits were made using both the Optomec LENS machine at The Ohio
State University, Columbus, OH, and the laser direct manufacturing facility at
Lockheed Martin Aeronautics Co., Fort Worth, TX. The system at Lockheed Martin
is essentially the same as the LENS system described in Chapter 1 with a few key
differences. The Lockheed Martin system does not have an oxygen purification
system, so the content of the atmosphere gas is not precisely known. Deposition
was usually conducted after three or four cycles of pumping down to one torr and
backfilling with argon to atmospheric pressure. The samples made on the LENS
system were monitored for oxygen content which was typically kept below 5 ppm.
Process parameters for all of the laser deposited samples are shown in Table 4.1.
44
V'o
..J::>.
0.015"
12"/min
20"/min
12"/min
12"/min
TRAVEL
SPEED
300W
425W
300W
300W
POWER
Table 4.1: Process parameters for the various deposits. The CP-Ti and CP-Ti + B samples were made at OSU and
the Ti-8Al-1Er samples were made at Lockheed Martin.
Ti-8Al-1Er (thin)
NIA
0.010"
0.010"
0.500" x 1.000" x 0.500"
Ti-8Al-1Er (bulk)
1.000" radius x 0.500"s
0.010"
0.015"
0.625" x 0.500" x 0.250"
CP-Ti + lB
0.010"
0.015"
0.625" x 0.500" x 0.250"
CP-Ti
LAYER
HEIGHT
ROW
SPACING
DIMENSIONS
SAMPLE
The bulk samples were made by building a crosshatch structure. The travel direction for each successive layer rotated 90°. The thin wall deposits were made by
depositing a helical structure one laser pass wide.
All of the bulk samples were cross-sectioned for microstructural evaluation.
Part of the sample was mounted in conductive resin, part was core drilled to make
TEM specimens, and part was kept for heat treatment. The bulk Ti-8Al-1Er samples
were taken from two different deposits made on two different days at the facility at
Lockheed Martin. Because of this, different levels of interstitial elements were
expected which can affect the microstructure and properties. The thin-walled Ti8Al-1Er samples, also made at Lockheed Martin, were cut from an eight-sided
helical structure and the longitudinal sides were mounted in conductive resin. The
CP-Ti and CP-Ti + B samples were deposited with the LENS unit at Ohio State.
These samples were deposited on different days as well, but because of the accurate
atmosphere monitoring, little difference was expected in their interstitial contents.
Various heat treatments were done on some of the samples to precipitate out
any supersaturated elements in solid solution. The heat treatment schedule for the
various deposits is shown in Table 4.2. Each sample was encapsulated in quartz
tubing, evacuated, and backfilled with 0.2 atmosphere argon to prevent oxidation
during heat treatment. The samples were heat treated in a tube furnace and allowed
to air cool.
Scanning electron microscope (SEM) samples were prepared by mounting in
conductive resin. These samples were ground flat using various grit SiC metallographic paper, then polished with 1.0 µm diamond paste. Final polishing was done
using 0.05 µm colloidal silica. These samples were then etched using Kroll 's reagent: 2 mL HF, 4 mL HN03 , and water to 1000 mL. Preparation of the as-received
powder was done using cold mount epoxy resin. The liquid epoxy was mixed with
46
SAMPLE
TIME
CP-Ti
2 hrs
100°c
CP-Ti+ lB
2 hrs
100°c
Ti-8Al-1Er (bulk and thin)
30min
560°C
Ti-8Al-1Er (bulk and thin)
2 hrs
100°c
Ti-8Al-1Er (bulk and thin)
100 hrs
100°c
TEMPERATURE
Table 4.2: Heat treatment schedule for the various as-deposited samples.
47
the hardener and a small amount of powder was stirred in. The mixture was then
poured into a plastic mount and allowed to harden overnight. The specimen surface
was prepared similarly to the other samples, however, the finished specimen had to
be carbon sputter coated to provide a conductive path for analysis in the SEM.
Transmission electron microscope (TEM) specimens were prepared by
electro-discharge machining (EDM) the deposits using a core drilling attachment.
The resulting 3 mm cores were then sliced into wafers 300 µm thick using the wire
EDM attachment. The samples were then ground using 240 grit SiC paper to a
thickness of about 100 µm. Some of the samples were then dimple-ground to a
thickness of 25 µm and milled using a Gatan ion mill. Other samples were prepared
by jet electropolishing. The 100 µm disks were thinned using a Struers Tenupol-3
twin jet polisher with voltage of 12 V and current of 10 mA. The polishing solution,
a mixture of 600 mL methanol, 350 mL butyl alcohol, and 60 mL perchloric acid,
was maintained at a temperature of -40°C to -60°C using liquid nitrogen. Some
samples that were not successfully thinned in the jet polisher were subsequently ion
milled to achieve an acceptable thin area.
X-ray diffraction was performed on a Scintag PAD-V x-ray diffraction unit
using CuKa radiation. The samples were prepared by grinding the oxide off the
surface of the deposits using SiC paper. Scans for the CP-Ti and CP-Ti + B were
done using a 28 range of25 to 50 degrees at a scan rate of 0.4 degrees per minute.
The Ti-8Al-1Er scans were done using a 28 range of20 to 110 degrees with a scan
rate of one degree per minute. Electron microscopic analysis was performed on a
Philips XL-30 FEG SEM and a Philips CM200 TEM. Both microscopes had EDAX
EDS attachments for doing spectroscopic microanalysis. Seven hardness measurements for each sample were taken with a Buehler Micromet II Vickers
microhardness indenter with a 500 g load and a 10 second dwell time.
48
CHAPTERS
RAPID SOLIDIFICATION
Rapid solidification is an inherent characteristic of the Laser Engineered Net
Shaping (LENS) process. The term "rapid solidification," however, is open to argument
in that some clarification is needed to determine a critical solidification rate that qualifies
as "rapid." Previous work with rapidly solidified dispersion hardened titanium alloys
quotes solidification rates on the order of 103 to 106 K/sec [ 1]. An estimation for the
solidification rate in LENS materials is crucial in determining whether or not this process
will succeed in creating a bulk, dispersion hardened material. This chapter combines
experimental results and theory to evaluate the solidification behavior in LENS material..
A heat transfer model developed for welding processes is also investigated to determine
its applicability to the LENS process.
Heat transfer in welding processes was fist examined by Rosenthal [2] and later
by Adams [3]. They made calculations based on a point heat source for an infinite threedimensional workpiece and a line source for a thin, semi-infinite flat workpiece. If these
equations are valid for laser deposited materials with complex geometry, then solidification rates can be calculated for any material at any specified laser power and traverse
speed. These equations must be slightly modified, however, to account for the unique
characteristics of the deposition process. The LENS process differs from typical high
energy density welding processes in that the two-dimensional, flat workpiece can now be
considered parallel to the laser rather than normal to it (i.e. the thin-wall deposit becomes
49
-~
t
t
v
Figure 5.1: Schematic of isotherm in thin plate with line heat source showing the
adaptation made to account for laser processing conditions.
the two-dimensional, flat workpiece). The radial distance, r, is directed down into the
deposited material, indicating the melt pool depth into the workpiece (Figure 5 .1 ). A
factor of 1/2 must then be inserted into the line source equations to account for halving
the isotherm volume. Since the deposit is very thin (<1 mm), a line source is still assumed and two-dimensional heat transfer is applicable.
Assumptions are necessary to simplify the calculations. The latent heat of fusion
is ignored even though the solidification rate inside the fusion zone is considered. The
effects of injecting powdered material into the weld pool are also ignored. Rosenthal
claims his solution is not valid for heat flow in the fusion zone. This is because in his
calculations the temperature increases to infinity at the location of the point source which,
of course, is not possible. Since solidification rates through the freezing range are of
interest, temperatures high above the melting temperature are ignored and Rosenthal's
solution is assumed to be valid. Simple calculations of the peak temperatures at various
50
distances from the heat source can determine the edge of the melt pool. Rosenthal's
solution for the two-dimensional line source problem is
T-T =-q-exp(Pez)K0 (~Pez 2 +Per 2 )
00
2nkt
5.1
where Pe is the Peclet number defined by
_ pCPVz
Pez -
2
k
_ pCPVr
and Per -
2
k
and
= thermal conductivity
= density
= specific heat
p
v = velocity of heat source
q = heat input from source
t = thickness
T = temperature at point (z, r)
Tcc = temperature of substrate at furthest distance from deposit
K n = Bessel function of second kind and order zero
k
r
c
For validation of these calculations, 316 stainless steel powder was deposited and analyzed. Values of k, p, and Cp for 316 stainless steel are 16.442 W /m°C, 8,027 .17 kg/m 3,
and 502.416 J/kg°C respectively [4]. The LENS process requires approximately 280 W
input power to produce a stainless steel deposit one millimeter thick. An approximation
for the Bessel function, K 0 , by Adams states that for P > 2
K0 (P);:::: exp(-P)
/n
5.2
~2P
and by taking (8Pe/8Pe) = 0 to find the maxima, the peak temperature at any distance
from the heat source path can be determined with
T -T::::
00
P
-
q
kt( Per ),,}8ne
5.3
where e = 2.7183. Using Equation 5.3, the calculated distance from the heat source to the
approximate location of the liquid-solid interface is about 0.5 millimeter. Thus, the
distance from the heat source is relatively large and Equation 5.1 can be applied.
51
Adams states that solidification rates do not vary significantly at small distances
from the weld center line. Hence, solidification rates on the weld center line are assumed
to be representative of the solidification rate through the entire radial distance, r, at a
linear distance, z, from the heat source. The solidification rates are then
-[dTJv
= 2nk v (r-r",)2
dz
q
5.4
for three-dimensional heat transfer and
dT
-[dz
2
Vt
JV= 2nkpCP q
(
3
)
5.5
(T-T
00 )
for two-dimensional heat transfer. These equations were used to plot the solidification
rate versus the thickness of deposit, as shown in Figure 5.2. This plot shows that the heat
transfer is two-dimensional up to about 2.5 mm with a transition zone between 2.5 and
3.2 mm and three-dimensional for thicknesses greater than 3.2 mm. The two-dimensional
equation was then used to calculate a solidification rate for the thin-wall sample through
the freezing range. The calculated rate is about 1600 K/sec for a one millimeter thick
deposit (q
=280 W, V= 0.01092 mis, T= 1658 K).
Equation 5.4 was used to calculate
the solidification rate in the bulk sample. The calculated rate is about 7500 K/sec with
the same deposition parameters as the thin-wall specimen (q
= 1658 K).
= 280 W, V = 0.01092 mis, T
It should be noted that these solidification rates are only valid when the
temperature at T is assumed to be room temperature. While this may be true for the first
00
few passes of the laser, it is not necessarily true for the entire deposition of the object.
The substrate begins to increase in temperature and affects the thermal gradient within the
deposit. This causes the solidification rate to decrease as deposition time proceeds.
Experimental determination of the solidification rates in LENS deposits was made
through microscopic evaluation of 316 stainless steel samples. The as-deposited samples
were cross-sectioned and the cell spacings for both bulk and thin-wall deposits were
52
20,000 .-------.------..,..-----.,...-------r---s------.
15,000
~
u
(!)
<Zl
~
'-"
(!)
......
('j
0:::
10,000
00
-u
.5
0
0
5,000
oi....--~=--...1.....
0
____
_j__ _ _ _-1..._ _ _ _---1._ _ _ ____J
2
3
4
5
Thickness (mm)
Figure 5.2: Graph showing the cooling rate versus thickness based on the modified
equations (Equations 5.4 and 5.5) for heat transfer in welding processes. The flow can be
considered two-dimensional up to about 2.5 mm where it transitions to mixed and threedimensional for any thickness over approximately 3.2 mm.
53
examined in a scanning electron microscope. These cell spacings can be empirically
related to the solidification rate by
A.= 25£-o.2s
5.6
where A, is the cell spacing in microns and£ is the solidification rate in K/sec [5]. Figures
5.3 and 5.4 show SEM micrographs of the as-deposited microstructures with arrows
indicating the features used to measure the cell spacing. Figure 5.3 shows a cross-section
of a bulk sample (0.5 in x 0.5 in x 0.3 75 in) deposited at 280 W, 26 in/min travel speed,
and a 0.015 in layer height. The measured spacing was approximately 2.2 µm, giving a
solidification rate of about 6000 K/sec. Figure 5.4 shows a cross-section of a thin-wall
deposit(~
1 mm thick) deposited at 280 W, 26 in/min travel speed, and 0.015 in layer
height. The measured cell spacing was approximately 1.9 µm, giving a solidification rate
of about 10,000 K/sec, which seems high for stainless steel based on previous experimental work [6]. These rates, however, were not constant throughout the entire sample. The
bulk sample had cell spacings as small as 1.05 µm (83,000 K/sec) and the thin-wall
sample had spacings as small as 1.41 µm (29,000 K/sec).
The measured results do not correlate well with the heat transfer theory developed
by Rosenthal and Adams. According to the theory, the thin-wall deposit should have a
lower solidification rate since thermal transfer is limited to two dimensions. Both sample
geometries, however, showed a wide range of overlapping solidification rates. There
could be a number of reasons for this behavior. Convective effects are totally ignored in
the model. While this may not be significant for bulk samples, it could play an important
role in thin-wall samples. Argon gas is used to carry the powder to the molten pool. This
argon blowing on the sample is likely to cause heat loss through forced convection.
Bonding of the deposit onto the substrate could also have an effect on the heat transfer
through the sample. Often the first few layers do not fuse together completely since the
substrate is cold and heat transfer is rapid. These first few layers are crucial in establish54
Figure 5.3: Scanning electron micrograph of laser dep<?sited 316 stainless steel sample.
Sample was taken through vertical cross-section of a bulk deposit 0.5 in x 0.5 in square
by 0.375 in tall. Arrows indicate cell spacings used to estimate cooling rates.
55
Figure 5.4: Scanning electron micrograph of laser deposited 316 stainless steel sample.
Image was taken through vertical cross-section of cylindrical thin-wall sample(i.e. single
laser pass). Arrows indicate cell spacings used in estimating cooling rates.
56
ing a contact for heat transfer from the sample into the plate. In bulk samples this could
create an insulating layer between the sample and the substrate, thus reducing the solidification rate. Also, internal porosity throughout the sample can create hot spots that interfere with heat transfer and slow the solidification rate in local areas.
The model given by Rosenthal and Adams is not very accurate when predicting
solidification rates in LENS processed material. The solidification rates in the deposits
vary considerably and the heat flow models are not appropriate for anything other than
rough, order of magnitude calculations. Better understanding and control of the LENS
process might someday produce deposits with a more constant cooling rate. For now, the
experimental results verify that the solidification rates are sufficient (103 to 104 K/sec) to
produce the desired dispersion strengthening effects in bulk titanium samples.
REFERENCES
1. Lewis, G. K., J. 0. Milewski, R. B. Nemec, D. J. Thoma, M. Barbe, and D. Cremers,
Technical Report LA-UR-95-2845.
2. Rosenthal, D., Welding Journal 20 (1941).
3. Adams, C. M., WeldingJournal37 (1958).
4. Alloy Digest 12 Engineering Alloys Digest, Upper Montclair, NJ (1980).
5. Katayama, S. and A. Matsunawa, Proc. ICALEO (1984).
6. Lienert, T. J., R. J. Grylls, and H. L. Fraser, unpublished research (1998).
57
CHAPTER6
LASER DEPOSITED TITANIUM MICROSTRUCTURES
6.1 Commercially Pure Titanium LENS Deposits
Deposits were made using commercially pure (CP) titanium to investigate
the role of rapid solidification on the microstructure of unalloyed titanium. The
deposits were made using a travel speed of 12 inches per minute and 300 W of
power on the Optomec LENS unit at Ohio State. The deposits were made in an
argon atmosphere with an oxygen concentration below 5 ppm to reduce the hardening effects of interstitial elements.
A backscattered electron SEM micrograph of the as-deposited CP-Ti structure is shown in Figure 6.1. The microstructure consists of acicular alpha titanium
with an alpha lath spacing on the order of 1-2 µm. In CP titanium, two hardening
mechanisms should be present: grain size strengthening and solid solution strengthening. TEM examination (Figure 6.2) shows the average lath spacing to be approximately 1.4 µm. This lath spacing can be used to calculate an approximate hardening
contribution based on the Hall-Petch grain size relationship. Typical values for r 0 ,
the matrix friction force, and k, the Hall-Petch coefficient for titanium, are 78.5 MPa
and 400 MPa·µm 0·5, respectively [1]. Inserting 1.4 µm for the grain size gives a
yield strength of approximately 415 MPa for the grain size contribution alone. It
must be noted, however, that the lath boundaries are likely low angle boundaries that
do not significantly contribute to grain size strengthening according to the Hall58
Figure 6.1: BSE SEM micrograph of as-deposited CP-Ti showing a fine, acicular
martensitic microstructure.
59
(a)
(b)
Figure 6.2: Brightfield TEM micrographs showing alpha lath boundaries in CP-Ti in the
as-deposited condition (a) and heat treated for two hours at 700°C condition (b).
60
Petch relationship. The largest strengthening contribution is most likely due to the
interstitial solute content.
Hardness values were measured for the as-deposited samples to give an
approximate indication of the strength of this as-deposited structure. An empirical
relationship has been established between the Vickers hardness of titanium and its
yield strength. Numerically, the hardness (typically measured in kg/mm 2) is roughly
equivalent to three times the yield strength [2]. Hardness values for the as-deposited
structures are shown in Table 6.1. The as-deposited hardness is about 227 kg/mm 2
which corresponds to an estimated yield strength of approximately 690 MPa.
The difference between the strength values (hardness value estimation minus
the grain size contribution) yields the contribution from solid solution strengthening.
Thus, the solid solution effects of carbon, nitrogen, and oxygen contribute a combined 275 MPa to the overall strength of the deposit. It should be noted, however,
that this is an overly simplistic view of the complex hardening that is occurring in
these deposits. There could also be effects attributable to strain hardening since the
deposit is exposed to high thermal strains during rapid solidification. This contribution, however, will be ignored since the successive heating of previous layers
a~
the
deposit is formed allows many of the dislocations to be annealed out. This results in
reduction, but not elimination, of the thermal stresses induced during deposition.
Establishing a baseline strength estimation will become important when more
complicated titanium alloy systems are considered.
Heat treatment was done on the CP titanium sample to see if any microstructural changes would occur. The sample was heated at 700°C for two hours in a
reduced pressure argon atmosphere. The resultant microstructure showed no
changes, as seen in the TEM micrograph comparison in Figure 6.2. Hardness values
for the heat treated sample were also taken and are listed in Table 6.1. Although, the
61
Vickers
Hardness
227
252
244
251
Standard
Deviation
9.4
12.0
3.1
3.9
Ti-8Al-1Er thin wall (as deposited)
Ti-8Al-1Er thin wall (30 min@ 540°C)
Ti-8Al-1Er thin wall (2hrs@ 700°C)
Ti-8Al-1Er thin wall (lOOhrs@ 700°C)
387
392
396
402
8.5
4.8
10.5
12.9
Ti-8Al-1Er bulk (as deposited)
Ti-8Al-1Er bulk (30 min@ 540°C)
Ti-8Al-1Er bulk (2hrs@ 700°C)
Ti-8Al-l Er bulk (1 OOhrs @ 700°C)
406
364
433
425
8.1
6.6
3.5
9.8
CPTi
CP Ti (2hrs @ 700°C)
CP Ti + B (as deposited)
CP Ti + B (2hrs @ 700°C)
Table 6.1: Vickers hardness values for the various laser deposited samples.
62
heat treated sample showed higher hardness values, the measurement scatter was
high as indicated by the large standard deviation. The heat treated sample would be
expected to be weaker since any strain hardening due to the thermal stresses from
rapid solidification would be annealed out. Grain growth could also become have an
effect, though not much is expected at the tested time and temperature. The difference in the numerical hardness values will be treated as a statistical anomaly, however, since the measurement errors nearly overlap.
6.2 CP-Ti Plus Elemental Boron LENS Deposits
Previous work has shown that small amounts of boron added to titanium is
an effective grain pinning agent, thus producing a refined microstructure. Work has
also shown that boron additions can produce a dispersion of second phase particles
that provide strengthening via the Orowan relationship. A blend of powder was
made using CP titanium and elemental boron powder. One weight percent boron
was added to the titanium powder and was mixed thoroughly before being loaded
into the powder feeder. Deposition was performed at a travel speed of 12 inches per
minute and 300 W power. The samples showed a fairly homogeneous structure with
primary alpha titanium and interdendritic titanium-boron compound. Figure 6.3
shows a backscattered electron (BSE) SEM micrograph of the as-deposited structure.
Determination of the second phase composition was important in establishing microstructural stability. The possible presence of titanium-boron compounds
was assessed using x-ray diffraction. Figure 6.4 shows an x-ray diffraction scan of
the as-deposited Ti-B alloy compared with that of the as-deposited CP-Ti alloy. The
presence ofTiB 2 is found, but no evidence ofTiB is seen. At such low boron concentrations, the compound TiB is the expected phase according to the equilibrium
63
Figure 6.3: BSE SEM micrograph of as-deposited CP-Ti + lB sample.
64
5000
0
~~~~~~~~~~~~~~~~~~~~
25
30
35
40
45
50
15,000 >--------------------------·
10,000
5,000
30
35
40
45
50
Figure 6.4: X-ray diffraction scans for the CP-Ti deposit (a) and CP-Ti + IB deposit (b).
Note the presence of the TiB 2 peak and the lack of a TiB peak (28 = 42°).
65
Figure 6.5: TEM micrograph of TiB 2 particle showing the diffraction pattern (inset) used
for identification.
66
phase diagram, which is not valid in RSP alloys. This is also contradictory to other
work done on RSP Ti-B alloys which showed that the second phase was TiB, not
TiBr Figure 6.5 shows the results of electron diffraction conducted on a large
second phase particle inside the matrix. This particle was large enough and in a thin
enough part of the sample (i.e. little interference from the surrounding matrix) to
allow a diffraction pattern to be obtained. Calculated interplanar spacings from this
pattern match the interplanar spacings found in the literature for TiB 2 •
Although the microstructure was quite homogeneous, there were some areas
where alloying did not completely occur, as well as areas where there was an overabundance of second phase particles. Figure 6.6 shows a BSE SEM micrograph of
one of such region where there was a large volume fraction of precipitate. This
image represents one of several regions where this type of precipitation occurred, all
of which surrounded large voids in the sample. This effect can likely be attributed to
the heat flow through the part and, specifically, to a lack of heat flow around the
void. These voids provide an insulating barrier inside the deposit that interrupts heat
transfer through the sample. Thus, around the hole, the matrix retains heat and is
slower to cool from the beta phase field into the alpha phase region. The reduced
cooling rate allows more time for solute diffusion in the more open body-centered
cubic crystal structure of beta titanium, creating a local region high in second phase
content. Chemical inhomogeniety cannot be ruled out, however, since this deposit
was made from a blend of elemental powders. Local variations in the boron concentration could also produce the microstructure seen in Figure 6.6.
The sample was heat treated to examine any microstructural changes that
might occur. The sample was heated in reduced pressure argon for two hours at
700°C. The resultant microstructure showed little or no changes compared with the
as-deposited structure. This result is consistent with what has been seen in other
67
(a)
'
(b)
Figure 6.6: BSE SEM micrograph of CP-Ti + lB deposit near a large void. Note the
higher concentration of second phase adjacent to the void.
68
(a)
-
(b)
Figure 6.7: Brightfield TEM micrographs of the heat treated CP-Ti + lB sample showing
rod-like TiB 2 particles
69
LENS deposits. Subsequent layers made during the deposition process essentially
heat treat the microstructure, thus leaving the as-deposited structure fairly stable.
Figure 6.7 shows the heat treated microstructure. It appears that perhaps more Ti-B
compounds have precipitated out of solution with little or no coarsening of the
existing particles. This claim, however, cannot be confirmed since the deposit was
not entirely homogeneous and a TEM sample represents a very small sample of the
microstructure.
Vickers hardness measurements were taken on the as-deposited and heat
treated Ti-IB samples. The results are shown in Table 6.1. The boron-containing
deposit had a hardness of about 17 kg/mm2 higher than that of the CP titanium
deposit. The heat treated sample was slightly harder, but not sufficiently to conclude
that more precipitation had occurred upon heat treatment, making the alloy stronger.
The discrepancy between the heat treated CP titanium sample and the boron deposits
can be attributed to statistical error, again due to the high standard deviation of the
heat treated CP titanium sample.
Deposition of CP titanium with small additions of boron show that it is
possible to produce a dispersion of second phase particles within the matrix. The
phase seen in the as-deposited sample was the non-equilibrium TiB 2 structure. This
phase remained after a two hour heat treatment at 700°C. These results show that a
non-equilibrium microstructure can be achieved using LENS to process titanium
alloys. This non-equilibrium structure is solid evidence that rapid solification is
occuring in the process. With this knowledge, a more effective dispersion hardened
alloy can be produced from prealloyed powder.
70
6.3 Prealloyed Ti-8Al-1Er Powder
Blending elemental powders is attractive when considering the expense to
produce variable composition alloys and novel alloys in the prealloyed form. These
alloys, however, tend to be less homogeneous when deposited using the LENS
process. Powder morphology, density, and other physical differences can contribute
to compositional differences at the molten pool. To avoid these problems, a
prealloyed powder was made by Crucible Research of the composition Ti-8Al-1Er.
The erbium is added with the expectation that it will dissolve into solid solution in
the as-deposited structure and precipitate out as erbium oxide upon heat treatment.
The aluminum is added for two reasons. First, titanium can dissolve a large amount
of aluminum into solid solution, increasing the strength. Also, aluminum is a strong
alpha phase stabilizer, which is very important for high temperature stabilization of
the microstructure. The a phase of titanium is hexagonal close-packed compared
with the body-centered cubic
J3 phase.
The close-packed structure has less free
volume than the body-centered cubic phase, slowing the diffusion kinetics for solute
atoms. The presence of the
J3 phase in this alloy would allow the microstructure to
coarsen rapidly, especially at the a/J3 interface, thus reducing the benefits gained
through dispersion hardening.
Analysis of the as-received powder showed a roughly spherical structure
with a substantial number of satellites. Though not all of the powder fell within the
size range specified by the manufacturer, most was within the desired -80/+200
mesh range. The powder also contained some internal porosity which can contribute
to the level of porosity in the deposited structures.
Figure 6.8 shows a BSE SEM cross-sectional micrograph of the as-received
powder. The individual grains/colonies are on the order of 10 µm and consist of
acicular martensitic a titanium. Although the light phase in between the grain
71
Figure 6.8: BSE SEM micrograph showing microstructure of as-received Ti-8Al- l Er
powder. Note the presence of elemental erbium at the grain boundaries.
72
boundaries is too small to accurately measure with EDS, it can be assumed to be
elemental erbium. Its high contrast in the micrograph suggests a high Z-number
element (Zr; = 22, ZEr = 68). Furthermore, erbium has a lower melting point than
titanium (l 497°C and 1668°C respectively), so it would be expected to solidify last
at the grain boundaries. This observation is important for two reasons. First, titanium can dissolve substantial amounts of oxygen into solid solution. Elemental
erbium is unstable in the presence of oxygen; it would prefer to exist as an oxide
(Er20J This oxide would not melt below the melting point of titanium and thus
would not wet the grain boundaries as seen in Fig. 6.8. Therefore, it appears that the
erbium present in the powder is elemental erbium and not an erbium compound.
Second, the appearance of erbium at the grain boundaries in the as-received powder
provides insight into the solidification rate of the powder. Although erbium has
virtually no solid solubility in a-titanium at room temperature, high solidification
rates can increase the solubility limit. The powder was processed using gas atomization, which typically has relatively high solidification rates. It might be expected
that the erbium would be dissolved into solution in the as-received powder. The fact
that a large amount of the erbium did not dissolve could be an indication that exceptionally high solidification rates are needed to completely dissolve the erbium in the
titanium matrix. This is of some concern since this powder is expected to dissolve
the erbium and precipitate out dispersoids in the laser deposited and heat treated
states. If the erbium cannot be dissolved in the matrix through laser deposition, then
the resultant structure is not unique at all, but virtually identical to a conventionally
processed alloy of the same composition.
Elemental erbium in the as-received powder could also indicate that the
concentration of erbium is too high even for RSP material. The composition for this
particular alloy was chosen for the following reasons: first, the erbium level was
73
chosen so that it would dissolve in solution and react with the available oxygen in
the titanium matrix when heat treated. The nominal oxygen concentration quoted
from the manufacturer was 0.14%. Based on this value and on the chemical formula
for erbium oxide, Er20 3, an atomic ratio of 2/3 erbium per one atom of oxygen was
used to calculate the amount of erbium needed to fully react with the nominal
oxygen present. This calculated to roughly one weight percent erbium. Previous
work by Konitzer (Chapter 3) showed that a slight excess of oxygen is needed to
prevent the coarsening of the erbium particles. Some oxygen will be picked up
during the deposition process so that a slight excess of oxygen will be present in the
as-deposited structures. Although multiple compositions could empirically determine the appropriate erbium concentration, one composition was chosen due to the
high cost of custom prealloyed titanium powders.
6.3.1 Single Pass Ti-8Al-1Er Beads
Single passes were made using the Ti-8Al-1Er powder to determine how the
laser processed material would behave. When depositing multiple pass layered
structures, the heat from the adjacent layers can induce microstructural changes in
the deposits. For this reason, a single pass was needed to isolate the solidification
characteristics of the deposits. The single pass deposits were made on a Ti-6Al-4V
substrate with a laser travel speed of 12 inches per minute and a laser output power
of 400 W. Backscattered electron SEM micrographs of the single pass structure are
shown in Figure 6.9. Overall, the single pass structures showed very little evidence
of a second phase, though there were some regions containing erbium/erbium
compounds. The top image in Figure 6.9 shows a finely dispersed second phase
contained within the grain structure while the bottom image shows particles on the
grain boundaries. These observations indicate that the erbium could exist in three
74
(a)
(b)
Figure 6.9: BSE SEM micrographs of a single bead deposit ofTi-8Al- 1Er. Note the
dispersed erbium bearing-particles in (a) and the particles lined up along the grain
boundaries in (b).
75
different forms in the microstructure. The first form is in solid solution in the
titanium matrix. The area fraction of second phase is much smaller in the single
pass than in the as-received powder, so it is likely that some of the erbium has been
dissolved in the matrix. Second, the phase dispersed throughout the individual
grains is an indication that some of the erbium is reacting with oxygen (oxygen
dissolved in the matrix or oxygen picked up from the atmosphere) and forming an
erbium-oxygen phase within the titanium matrix. Finally, the erbium at the grain
boundaries indicates that there could still be elemental erbium in the deposit, similar
to that which was seen in the as-received powder.
6.3.2 Ti-8Al-1Er LENS Deposits
Deposits were made using the LENS process with the prealloyed Ti-8Al-1Er
powder received from Crucible Research. Several deposits were made with different geometries under different processing conditions. All the deposits were made on
a Ti-6Al-4V substrate using the laser direct manufacturing facility at Lockheed
Martin Aeronautics Company, Fort Worth, TX. Based on the analysis of geometry
dependent cooling rates (Chapter 5) and the precipitation effects from cooling rates
seen in the Ti-B samples, both thin wall (two-dimensional heat flow) and bulk
(three-dimensional heat flow) samples were made. In addition to the two different
geometries, three different heat treatments were selected to investigate the precipitation behavior at elevated temperatures and prolonged times.
Figure 6.10 shows BSE SEM micrographs of the as-deposited Ti-8Al-1Er
samples, both for bulk (top) and thin wall (bottom) deposits. It can easily be seen
that the erbium has produced a very fine distribution of second phase particles
homogeneously distributed throughout the acicular a-titanium matrix. This is
clearly illustrated in the TEM micrograph in Figure 6.11. There is a observable
76
(a)
(b)
Figure 6.10: BSE SEM micrographs of the bulk (a) and thin wall (b) as-deposited Ti8Al-1 Er samples.
77
Figure 6.11: Brightfield TEM micrograph of the bulk as-deposited Ti-8Al- l Er structure.
78
Figure 6.12: Brightfield TEM micrograph of thin wall Ti-8Al-1Er deposit showing a
bimodal distribution of second phase particles.
79
Figure 6.13: BSE SEM micrograph ofTi-8Al-1Er deposit showing bands of precipitates
coinciding with bands of porosity.
80
difference, however, in the appearance of the bulk sample compared to the thin wall
sample. This difference is possibly a result of the varied cooling conditions in the
different sample geometries. The thin wall structure is expected to exhibit twodimensional cooling due to its wall thickness (<l mm). This can reduce the cooling
rate through the beta region with an effect similar to that seen near the voids in the
Ti-B deposits. The slower cooling rate increases the amount of time spent in the
beta region and allows increased diffusion of the solute atoms, especially at the a/j3
interface. The result is large dispersoids (100 - 500 nm) with regions in between of
smaller dispersoids (50 -100 nm) on the order of those seen in the bulk deposit (see
Figure 6.12). The lath size in the matrix is also considerably larger in the thin wall
sample further supporting the claim that the cooling rate is less than that in the bulk
sample.
Figure 6.13 shows a BSE SEM micrograph of the bulk deposit containing
bands with a higher volume fraction of precipitates. Careful examination shows that
small voids, likely due to entraped gas in the powder, are contained within these
bands. These voids create localized hot spots within the deposit. Again, this is
similar to what was seen in the Ti-B samples near the voids and can be explained by
the same argument. These combined results show how the heat transfer through the
deposit is crucial in obtaining the desired microstructure. Internal voids, poor
substrate bonding, sample geometry, and other characteristics affecting heat flow
can all be influential in the way dispersoids form within the microstructure.
Figures 6.14 through 6.16 show the bulk and thin wall Ti-8Al-1Er deposits in
the heat treated condition. Very little change occurs through the various stages of
heat treatment. Figure 6.14 shows the samples that were stress relief annealed at
560°C for 30 minutes. No microstructural changes were seen in the deposits. Hardness values for the thin wall structure were slightly higher than the as-deposited
81
(a)
(b)
/
Figure 6.14: BSE SEM micrographs of the bulk (a) and thin wall (b) Ti-8Al-1Er samples
heat treated at 560°C for 30 minutes.
82
(a)
I
(b)
Figure 6.15: BSE SEM micrographs of the bulk (a) and thin wall (b) Ti-8Al-1Er samples
heat treated at 700°C for two hours.
83
(a)
(b)
Figure 6.16: BSE SEM micrographs of the bulk (a) and thin wall (b) Ti-8Al-1Er samples
heat treated at 700°C for 100 hours.
84
(a)
(b)
Figure 6.17: BSE SEM micrographs comparing particle size in the bulk as-deposited (a)
Ti-8Al- l Er and the 100 hour/700°C heat treated sample (b).
85
(a)
(b)
I
Figure 6.18: BSE SEM micrographs comparing particle size in the thin wall as-deposited
(a) Ti-8Al-1Er and the 100 hour/700°C heat treated sample (b).
86
values, but significantly lower in the bulk specimen (see Table 6.1 ). It is expected
that this type of heat treatment would reduce the number of dislocations compared
with the thermally stressed as-deposited structure. The differences in hardness are
most likely caused by local variations in the precipitate size and distribution, i.e.
sampling inconsistencies. The other heat treatments, 700°C for two hours and
700°C for 100 hours, did show some changes in the microstructure, as seen in
Figure 6.15 and Figure 6.16. The heat treatment mainly affected the matrix, resulting in very little particle coarsening. The alpha structure of the matrix coarsened in
the 100 hour heat treatment. The size and distribution of the erbium-bearing particles did not change much over the various heat treatments in the bulk deposits.
This is evidence that very little erbium is remaining in solution in the as-deposited
condition. Figures 6.17 and 6.18 show high magnification BSE SEM micrographs
of the as-deposited samples (thin wall and bulk) and the 100 hour heat treated
samples (thin wall and bulk). By attributing the particle distribution to sampling
error and noting the particle sizes, one can conclude that there is little or no coarsening of the erbium-bearing particles.
The thin wall structures showed a definite organized arrangement of particles
that became more apparent after the longer heat treatments. Close inspection of this
phenomenon in the TEM shows that the particles are lined up along grain boundaries
and dislocations (Figure 6.19). This behavior is expected since the erbium atoms are
larger than the titanium and would segregate to the more open areas of the lattice,
such as edge dislocations and grain boundaries. The lack of particle alignment in the
bulk samples can also be attributed to differences in the cooling rates. The thin wall
samples have longer dwell times in the beta region, thus forming the large scale
precipitates. This high temperature dwell also allows the smaller particles to segre-
87
(a)
•
."
(b)
Figure 6.19: Brightfield TEM micrographs showing how the precipitate particles line up
along lath boundaries and dislocations.
88
(a)
.
..
.,,
•
.
•
•
,,
..
t'
#
' ·,
.
...
•
.
.·.·,..
•
•
•
.
.
,
••• •• •
'.
.~ I •
,•.
•: .
•
•
(b)
., ..• •
·~
•
.
•
••
•
••
Figure 6.20: Brightfield TEM micrographs comparing the precipitate size and
distribution in the 100 hour heat treated bulk (a) and thin wall (b) Ti-8Al-1Er deposits.
89
gate to the grain boundaries and dislocations and to be precipitated out after successive laser passes.
The formation of larger particles in the thin wall specimens depletes the
matrix of erbium available for precipitation during heat treatment. Figure 6.20
shows TEM micrographs of the 100 hour heat treated bulk and thin wall samples.
Evidence of precipitation of second phase particles in the bulk sample is seen between the larger particles that were evident in the as-deposited condition. These
small precipitates are not present in the thin wall specimen due to the lack of erbium
in the matrix. Therefore, the lack of erbium left in the matrix after deposition of the
thin wall structures prevents futher precipitation of the second phase after long term
heat treatment.
Positive identification of the second phase particles was difficult due to their
small size and low volume fraction. An x-ray diffraction scan was performed on one
of the Ti-8Al-1Er samples and was compared to CP titanium. The results did not
provide sufficient evidence to positively identify the particles. Electron diffraction
in the TEM was also done on the Ti-8Al-1Er samples for phase identification. It
was difficult to obtain a clear pattern because the particles were surrounded by
matrix material that also diffracted the beam. An extensive search of the thinned
area in one of the TEM samples revealed a particle on the edge with little matrix
material surrounding it. By reducing the spot size of the beam and choosing the
smallest selected area aperture, a diffraction pattern of the sample was obtained with
little interference from the matrix. The sample was tilted to align the particle along
its [T 12] zone so that a selected area diffraction pattern could be taken. This pattern
was then compared with that of the [112] zone ofEr20 3 [3] obtained from computer
simulation. A bright field TEM micrograph of the particle is shown in Figure 6.21
along with the experimental and simulated diffraction patterns. The diffraction
90
(a)
(b)
(c)
Figure 6.21: Brightfield TEM micrograph of isolated precipitate particle (a) with
simulated (b) and experimental (c) [T 12] zone axis di ffration pattern.
91
• •1
I
••
••••
-·
Figure 6.22: Brightfield TEM micrograph showing preferred orientation particle
coarsening in the 100 hour heat treated Ti-8Al-1Er deposit.
92
Si
s
Mn
c
Cr
Ni
Cu
Sn
Fe
Al
Ti
Er
0
N
as
de osited
0.016
0.004
0.004
0.173
0.009
0.023
0.007
0.002
0.146
11.081
87.077
0.084
0.762
0.613
30min@
540°C
0.016
0.004
0.004
0.222
0.017
0.023
0.014
0.002
0.129
11.623
86.735
0.086
0.705
0.419
2hrs@
100°c
0.016
0.004
0.004
0.190
0.009
0.023
0.007
0.002
0.129
12.221
85.831
0.083
0.841
0.641
100 hrs@
100°c
0.016
0.003
0.004
0.150
0.009
0.023
0.007
0.002
0.081
11.242
87.232
0.081
0.706
0.387
Table 6.2: Chemical analysis (atomic percent) of bulk Ti-8Al-1Er deposits.
93
pattern reveals that the second phase has a cubic crystal structure and an arrangement of diffraction maxima that closely match that of the simulated pattern. From
this analysis, it can be concluded with a high degree of certainty that the second
phase particles are indeed Er20x. These particles might also contain carbon and/or
nitrogen, however, the structure is consistent with Er20 3• Although an attempt to
identify the orientation relationship was not made, it appears that a relationship does
exist. Figure 6.22 shows the microstructure in the 100 hours at 700°C condition.
The particles appear to be coarsening and their morphology is changing from spherical to rod-like. This elongation occurs in preferred directions within each individual
grain, evidence that an orientation relationship between the particles and the matrix
does exist.
Strengthening of Ti-8Al-l Er is achieved through a combination of effects.
First, the eight percent aluminum contributes through solid solution strengthening.
Contrary to the work done by Sastry et al [Chapter 3] showing that the RSP alloy Ti8Al-2Er had a 2 (Ti 3Al) precipitates present, no aluminum compound was found in
this alloy. Chemical analysis (see Table 6.2) of the bulk samples showed that the
average aluminum composition was 11.5 at% (6.9 wt%). Correlating this value
with the graph shown in Figure 2.3, a total yield strength of approximately 758 MPa
results from solid solution effects of aluminum. Interstitial solid solution strengthening was also achieved by the presence of carbon, nitrogen, and oxygen in solution.
Estimation of the total yield strength of this alloy was calculated based on
the equations presented in Chapter 2 for solid solution strengthening (Equation 2.1 ),
Hall-Petch strengthening (Equation 2.4), and Orowan strengthening (Equation 2.17).
Table 6.3 shows the values of the variables used, their references, and values for the
calculated quantities. Only the bulk as-deposited sample was used in the strength
94
Variable
µ
b
ro
v
&o
&c
&N
To
k
Calculated
d
A.
D
Value
45.6 GPa
0.295 run
l.18run
0.36
0.0485
0.0765
0.1525
78.5 MPa
400 MPav'µm
Reference
5
5
5
5
4
4
4
1
1
Value
60run
200 run
1.4 µm
Table 6.3: Table of experimental and calculated values used in strength calculations for
Ti-8Al-1Er deposits.
95
calculations since comparison between the estimated sample strengths would likely
be inaccurate.
Solid solution strengthening effects were calculated based on the chemical
analysis done on the bulk samples. One and a half times the erbium concentration
was subtracted from the oxygen value to account for the amount of oxygen tied up
in the dispersion compound Er20 3• The atomic percent of carbon, nitrogen, and
oxygen were multiplied by their average atomic misfit parameters [4] to give a
composite interstitial value. This value was inserted into Equation 2.1, resulting in
326 MPa of strengthening contribution from the interstitial content. Comparing this
value with the sum of the experimental values in Figure 2.4 shows a poor correlation. The combined interstitial effect is 700 MPa and subtracting off the value for
the friction force of 78.5 MPa leaves 622 MPa for the total combined effect. Thus,
the empirical results predict a strengthening value 91 % higher than that calculated
using the equations. Given this considerable discrepancy, an estimate of the interstitial strength contribution can not be accurately determined.
The Hall-Petch strengthening contribution was calculated based on the
experimentally measured lath spacing. As seen in the CP-Ti deposits, the lath
spacing in these alloys is also approximately 1.4 µm. Using a value of 400 MPa for
k, the Hall-Petch slope, a value of 338 MPa was calculated as the grain size strength-
ening contribution. Again, it is noted that these lath boundaries may be low angle
boundaries that do not effectively inhibit dislocation motion.
Particle-dislocation interaction is very apparent in the as-deposited microstructure as seen in Figure 6.23. To calculate the Orowan contribution to the
strength of the alloy, the parameters A., the mean interparticle spacing, and d, the
mean particle diameter, must be measured. SEM analysis can estimate the number
of precipitates per unit area, giving an approximate measurement of the number of
96
Figure 6.23: Brightfield TEM micrograph showing dislocation/particle interaction. Note
that the particles are unshearable (hard Orowan barriers).
97
barriers on any given slip plane. Equation 2.18 is then used to calculate an effective
center-to-center spacing. TEM analysis provides an average value for the particle
diameters. The mean interparticle spacing is then calculated by subtracting the
average diameter from the center-to-center spacing. The measurements on the bulk
as-deposited structure resulted in an average particle radius of 60 nm and an interparticle spacing of 80 nm (200 nm center-to-center spacing). Inserting these values
into the equation gives an Orowan contribution of 106 MPa. It must be noted that
significant error is possible when using SEM micrographs to calculate the number of
particles per unit area. Backscattered electrons can come from below the surface of
the sample and the high accelerating voltage used (25 Ke V) creates a large interaction volume. The result is that some particles from just below the surface appear to
be on the surface, thus giving a higher number of barriers per unit area of slip plane.
For lack of an acceptable equation for calculating the substitution solid
solution effect from the aluminum in the matrix, the graph in Figure 2.3 was used as
an estimate. The strengthening effect (minus the strength at zero solute concentration) was 379 MPa.
Summing these values gives a total estimate for the strength of the alloy.
This total value, plus the value for the matrix friction stress, 78.5 MPa, is approximately 1230 MPa. Note that the calculated value of the interstitial solid solution
effect was used (326 MPa), not the value from Figure 2.4. This estimate seems
rather high since there are no alpha phase titanium alloys with this kind of strength.
Hardness values, however, indicate that this high strength value might actually be a
good estimation. The empirical relationship between Vickers hardness and the yield
strength was shown before to be HV =3Y. Performing this calculation on the
hardness values for the deposits show that the predicted strength, based on a hardness value of 400 kg/mm2 , is approximately 1280 MPa. Thus, both the calculated
98
and experimental evidence indicate that the Ti-8Al- l Er alloy may have a yield
strength on the order of 1240 MPa.
6.4 Conclusions
Laser engineered net shaping (LENS) technology was used to create rapidly
solidified titanium alloys. These alloys consisted of commercially pure (CP) titanium, CP-Ti + lB, and Ti-8Al-1Er. The CP titanium showed a martensitic alpha
titanium structure with a refined grain size on the order of 1.4 µm. Heat treatment at
700°C for two hours produced little change. CP-Ti + 1B was deposited from elemental blends of powders. The boron alloyed with the titanium in-situ and formed
a fine dispersion ofTiB 2 particles. Heat treatment at 700°C for two hours produced
little change in the microstructure. Both thin wall and bulk samples were made of
the Ti-8Al-1Er alloy. The bulk samples showed a fine distribution of spherical Er20 3
particles on the order of 60 nm with a mean interparticle spacing of about 80 nm.
The thin walled specimens had a bimodal distribution of large particles (200-500
nm) and small particles {<100 nm). This difference in microstructure is attributed to
the differences in cooling rates between the thin wall and bulk geometries. Cooling
rates through the beta region dictate the size and distribution of particles in the
matrix. Defects such as voids cause local heat flow differences and can effect
change on microstructural development. Control of the heat flow through the
sample is critical in achieving the desired microstructure and properties. Strength
estimates, based on microhardness values and calculations, show that these alloys
can exhibit high room temperature strength and can retain these strengths even after
long term high temperature exposure.
99
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