PHASE TRANSFORMATIONS AND EQUILIBRIA OF

Transcription

PHASE TRANSFORMATIONS AND EQUILIBRIA OF TITANIUM PLATINUM ALLOYS IN
THE COMPOSITION RANGE 30-50 ATOMIC PERCENT PLATINUM
by
Karem E. Tello Araya
UMI Number: 3546038
All rights reserved
INFORMATION TO ALL USERS
The quality of this reproduction is dependent upon the quality of the copy submitted.
In the unlikely event that the author did not send a complete manuscript
and there are missing pages, these will be noted. Also, if material had to be removed,
a note will indicate the deletion.
UMI 3546038
Published by ProQuest LLC (2012). Copyright in the Dissertation held by the Author.
Microform Edition © ProQuest LLC.
All rights reserved. This work is protected against
unauthorized copying under Title 17, United States Code
ProQuest LLC.
789 East Eisenhower Parkway
P.O. Box 1346
Ann Arbor, MI 48106 - 1346
A thesis submitted to the Faculty and Board of Trustees of the Colorado School of Mines in
partial fulfillment of the requirements for the degree of Doctor of Philosophy (Materials Science).
Golden, Colorado
Date:
Signed:
Karem E. Tello Araya
Signed:
Dr. Michael J. Kaufman
Thesis Advisor
Golden, Colorado
Date:
Signed:
Dr. Michael J. Kaufman
Professor and Head
Department of Metallurgy and Materials
ii
ABSTRACT
The Ti-Pt phase binary phase diagram and the corresponding phase transformations in the
composition range 30-50 at.% Pt have been investigated using a variety of characterization methods (DTA, SEM and TEM). This study was inspired by ongoing work on some experimental Ti-NiPt and Ti-Pt-Ni-Hf high temperature shape memory alloys that were found to contain unexpected
phases not reported previously in such alloys. Furthermore, close analysis of the peritectoid invariant proposed by Biggs et al. revealed a range of confusing and somewhat contradictory results and,
as a result, it was decided to attempt to determine the true nature of the diagram in this composition range and to understand the complicating effects of interstitial contamination on the observed
microstructures and phase equilibria.
The microstructure of as-cast and heat treated alloys contains more than two phases after
equilibration treatments suggesting interstitial contamination. In addition, the microstructures revealed that the peritectoid transformation (Ti3 Pt+β-TiPtTi4 Pt3 ) proposed in the literature exists
but, because of sluggish transformation kinetics, the actual peritectoid reaction is limited and does
not account for the observed DTA peaks that Biggs et al. used to estimate the invariant temperature. Rather, it will be shown that the peaks are due to the transformation of β-TiPt to a lamellar
β-TiPt+Ti4 Pt3 structure at approximately 1230 ◦C. In addition, a modification to the phase diagram
is proposed based on other experimental evidence.
Characterization of the various phases observed in the microstructures (using SADP and
CBED in the TEM) confirmed the presence of the known phases Ti3 Pt and α-TiPt. In addition, a
new phase with stoichiometry Ti5 Pt3 was observed in both as-cast and heat treated samples. This
phase is shown to be stabilized by oxygen and to have a hexagonal structure with lattice parameters
a ∼ 8.0 nm and c ∼ 5.0 nm (space group P63 /mcm). The Ti4 Pt3 phase appears to be a true binary
phase that tends to be highly faulted and be structurally related to the Ti5 Pt3 phase with a pseudohexagonal structure with a ∼ 7.96 nm and c ∼ 23.6 nm. Detailed electron diffraction evidence
iii
indicates that the crystal structure is probably triclinic although it was difficult to determine the
actual point and space group.
iv
TABLE OF CONTENTS
ABSTRACT . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .
iii
LIST OF FIGURES . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .
vii
LIST OF TABLES . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .
xiii
ACKNOWLEDGMENTS . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .
xiv
CHAPTER 1
INTRODUCTION . . . . . . . . . . . . . . . . . . . . . . . . . . . .
1
CHAPTER 2
BACKGROUND AND LITERATURE REVIEW . . . . . . . . . . . . .
6
2.1
Ti-Pt phase diagram . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .
2.2
Shape Memory Effect (SME) and Superelastic Effect (SE) . . . . . . . . . . . . . 10
2.3
8
2.2.1
Thermodynamics of the Thermoelastic Martensitic Transformation . . . . 13
2.2.2
Shape Memory Effect . . . . . . . . . . . . . . . . . . . . . . . . . . . . 15
2.2.3
Superelastic Effect (SE) . . . . . . . . . . . . . . . . . . . . . . . . . . . 15
Organization of the thesis . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 18
CHAPTER 3
CHARACTERIZATION OF PHASES . . . . . . . . . . . . . . . . . .
19
3.1
Experimental Methodology . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 19
3.2
Results . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 20
3.2.1
3.2.2
CHAPTER 4
4.1
Ti-34Pt Alloy . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 20
3.2.1.1
Faulted Phase: Ti4 Pt3 . . . . . . . . . . . . . . . . . . . . . . . 21
3.2.1.2
Ti5 Pt3 Phase . . . . . . . . . . . . . . . . . . . . . . . . . . . . 26
3.2.1.3
Globular Ti4 Pt(C,N) Phase . . . . . . . . . . . . . . . . . . . . 31
Discussion . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 33
Ti-Pt PHASE DIAGRAM . . . . . . . . . . . . . . . . . . . . . . . . .
38
Experimental Methodology . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 38
v
4.2
4.3
4.4
4.5
4.1.1
Differential Thermal Analysis (DTA) and Heat Treatment . . . . . . . . . 38
4.1.2
Scanning Electron Microscopy (SEM) and Energy Dispersive Spectroscopy
(EDS) . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 42
4.1.3
Transmission Electron Microscopy (TEM) . . . . . . . . . . . . . . . . . 43
As-cast Microstructures . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 43
4.2.1
Hypoeutectic Alloy: Ti-31Pt . . . . . . . . . . . . . . . . . . . . . . . . . 43
4.2.2
Eutectic Alloy: Ti-35Pt . . . . . . . . . . . . . . . . . . . . . . . . . . . . 44
4.2.3
As-Cast Ti-39Pt Alloy . . . . . . . . . . . . . . . . . . . . . . . . . . . . 48
4.2.4
4.2.5
Microstructures After Heat Treatment . . . . . . . . . . . . . . . . . . . . . . . . 54
4.3.1
Ti-31Pt Alloy . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 54
4.3.2
Ti-35Pt Alloy . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 57
4.3.3
Ti-39Pt Alloy . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 58
4.3.4
Ti-44Pt Alloy . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 61
Discussion . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 63
4.4.1
Transformations Observed in the as-cast Alloys . . . . . . . . . . . . . . . 63
4.4.2
Transformation β-TiPt β-TiPt+Ti4 Pt3 . . . . . . . . . . . . . . . . . . . 72
4.4.3
Transformation of Ti4 Pt3 from α-TiPt at low temperatures . . . . . . . . . 77
4.4.4
Microstructure of the Ti-44Pt alloys after heat treatment . . . . . . . . . . 82
Sequence of Transformations and Summary of Observed Phases . . . . . . . . . . 84
CHAPTER 5
CONCLUSIONS AND FUTURE WORK . . . . . . . . . . . . . . . .
96
REFERENCES CITED . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 100
APPENDIX A
XRD AND NEUTRON DIFFRACTION RESULTS . . . . . . . . . . . 104
A-1 Methodology . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 104
A-2 Results . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 106
APPENDIX B
STRUCTURE FACTORS FOR XRD AND NEUTRON DIFFRACTION
vi
109
LIST OF FIGURES
1.1
Quaternary as-cast alloy 51.8Ti-34.4Pt-6.5Ni-6.3Hf (at.%) (a) BEI (b) TEM image revealed the presence of a highly faulted phase (F) and martensite (M) in its
microstructure. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .
2
1.2
Ti-Pt phase diagram proposed by Biggs et al. (from [1]). . . . . . . . . . . . . . .
2
1.3
SEM results reported by Biggs et al.. (a) Backscattered Electron Image (BEI) of
Ti-37.6Pt as-cast alloy showing dendrites of Ti4 Pt3 with Ti3 Pt+Ti4 Pt3 eutectic in
the interdendritic regions. (b) BEI of Ti-42.7Pt as-cast alloy. The authors observed
that the dendrites were comprised of cores of α-TiPt phase and Ti4 Pt3 at the peripheries . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .
3
Alloying element effect on the martensitic start and finish transformation temperatures for the binary alloys TiAu, TiPd, and TiPt . . . . . . . . . . . . . . . . . . .
7
Transformation temperature versus percentage of alloying element added to Ti50at.%Ni alloy . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .
7
2.1
2.2
2.3
Phase Diagram of Ti-Pt system proposed by: (a) Nishimura and Hiramatsu in
1957 [2], (b) Murray in 1982 [3], (c) Biggs et al. in 2004, and (d) Li et al. in
2008 . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 10
2.4
Schematic of the martensitic transformation (MT) showing the characteristic temperatures involved during the transformation . . . . . . . . . . . . . . . . . . . . . 11
2.5
Structural relationship between the parent B2 phase (TiPt-β) and the martensite
B19 (TiPt-α) . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 12
2.6
Schematic representation of the free energy curves for the austenite and martensite
phases. ΔTF is the undercooling required for the martensitic transformation in
the forward direction, while ΔTR is the superheating required for the martensitic
transformation in the reverse direction. . . . . . . . . . . . . . . . . . . . . . . . . 14
2.7
(a) Schematic of the twinned martensite which transforms upon cooling in the absence of an applied load, while (b) shows a schematic of the detwinned martensite
which forms at low temperatures when a load sufficiently high is applied to reorient
the martensitic variants into a preferred orientation variant . . . . . . . . . . . . . 16
2.8
The austenite phase existing at a temperature above Af is subjected to a stress
above σMf resulting in fully detwinned martensite. When the specimen is unloaded
the austenite phase is fully recovered. This is known as superelastic effect . . . . . 17
vii
3.1
Low magnification (a) and high magnification (b) BEIs of the Ti-34Pt as-cast alloy
showing the primary dendrites of Ti3 Pt (dark phase) and multi-phase interdendritic
eutectic comprised of Ti4 Pt3 and α-TiPt (which transforms from β-TiPt upon cooling). (c) Bright Field TEM (BFTEM) of the interdendritic regions of the Ti-34Pt
as-cast alloy. As can be seen there are multiple phases in these regions including
three that are actually new phases, namely Ti5 Pt3 , Ti4 Pt3 and the globular phase. . . 21
3.2
Convergent Beam Patterns (CBPs) and SADPs of the Ti3 Pt phase in Figure 3.1a:
(a) 011 CBP and (b) 111 CBP and (c) 001 SADP. These patterns are consistent with the cubic structure of Ti3 Pt (a = 0.503 nm). . . . . . . . . . . . . . . . . 22
3.3
SADPs from the faulted Ti4 Pt3 phase in Figure 3.1c. As discussed in the text, the
spacing of the spots in the unstreaked row and the spacings orthogonal to this row
are readily indexed using hexagonal indices with a ∼0.8 nm and c ∼ 2.36 nm
(included in the table). . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 24
3.4
BFTEM (a) and high resolution electron micrograph (HREM) (b) of a region of
the faulted phase in the Ti-34Pt as-cast alloy that appears to contain no faults. This
was rare as, typically, this phase is heavily faulted throughout. . . . . . . . . . . . 25
3.5
SADP of the [1120] (a) and [1100] (b) zones (referred to the hexagonal indices) of
an unfaulted region of the faulted phase revealing the lack of a mirror plane in the
zero order Laue zone. Thus, the symmetry of this pattern is either 2 or 1. . . . . . . 25
3.6
(a) 1100 CBP of an unfaulted region of the “faulted” pseudo-hexagonal phase
Ti4 Pt3 . While it appears that there might be two orthogonal mirrors, the lack of a
0001 mirror in the 1120 pattern (Figure 3.5a) and the observed Kikuchi bands at
high angles (marked) that are not mirrored to the bottom indicate that the whole
pattern symmetry is 1-fold. (b) SADP of the faulted phase taken near B= 0001.
As can be seen, the spot spacings and the angles between the systematic rows are
very close to those expected for a 3-fold or 6-fold axis. . . . . . . . . . . . . . . . 26
3.7
CBPs from the Ti5 Pt3 phase in the Ti-Pt alloys along (a) 0001, (b) 1120 and (c)
1100 directions. Note that the patterns in (a) are 6mm while those in (b) and (c)
have 2mm symmetry in the whole pattern consistent with the 6/mmm point group.
Also, note the dynamic absences (arrowheads) in (c). . . . . . . . . . . . . . . . . 27
3.7
Continued . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 28
3.8
SADPs from the Ti5 Pt3 phase in the Ti-Pt alloys along (a) 0001, (b) 1120 and
(c) 1100 directions used to determine the lattice parameters and extinctions of
this structure. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 29
3.9
Ti-42Pt alloys (a) 4 h at 1200 ◦C +WQ. (b) 2 h at 1260 ◦C +WQ. . . . . . . . . . . 30
3.10 (a) 001, (b) 111 and (c) 011 CBPs from the globular Ti4 Pt phase. The symmetries of these patterns are consistent with an m3m point group and the observed
forbidden reflections with an Fm3m space group. The weak odd reflections in the
011 SADP in (d) indicate a B1 (NaCl) ordering. . . . . . . . . . . . . . . . . . . 32
viii
3.11 (a) B = 0001 of both Ti5 Pt3 and Ti4 Pt3 phases, (d) B = 1123 of Ti5 Pt3 (left
side) and both phases (right side) and (e) B = 1122 of Ti5 Pt3 (left side) and both
phases (right side). . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 36
4.1
Differential thermal analysis results obtained for Ti-31Pt, 35Pt, 39Pt, and 44Pt alloys. 40
4.2
(a) Ti-Pt phase diagram showing the alloys heat treated at three different temperatures to bracket the peaks observed during the DTA experiments. (b) Schematic of
the heat treatments performed on the alloys. . . . . . . . . . . . . . . . . . . . . . 41
4.3
Schematic of the location where the samples for metallographic and microstructural analysis were obtained from the melted buttons. . . . . . . . . . . . . . . . . 42
4.4
Ti-31Pt alloy in the as-cast condition. (a) 1000X (b) 28000X (c) EDS Spectra . . . 44
4.4
Continued . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 45
4.5
Ti-35Pt alloy in the as-cast condition. (a) 5,000X. (b) TEM image . . . . . . . . . 46
4.5
Continued . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 47
4.6
Ti-39Pt alloy in the as-cast condition. (a) BEI 1,000X (b) BEI 5,000X, (c-d) 45,000X 49
4.6
Continued . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 50
4.7
Ti-42Pt alloy in the as-cast condition. (a) 1,000X, (b) 5,000X and (c) 8,000X . . . 51
4.7
Continued . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 52
4.8
Ti-42Pt alloy in the as-cast condition. (a) 35,000X (b) 75,000X . . . . . . . . . . . 52
4.9
Ti-44Pt alloy. (a) 2,000X (b) 35,000X . . . . . . . . . . . . . . . . . . . . . . . . 53
4.9
Continued . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 54
4.10 Continued . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 56
4.11 Ti-35Pt alloy heat treated at (a) 1260 ◦C (HT1), (b) 1050 ◦C (HT2) and 800 ◦C (HT3). 57
4.11 Continued . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 58
4.12 (a) Ti-39Pt alloy heat treated at 1260 ◦C (HT1). The microstructure is composed of
dendrites of β-TiPt that later transform to α-TiPt upon quenching, Ti3 Pt and Ti5 Pt3 .
SADPs from Ti3 Pt, B=001, and Ti5 Pt, B=2113 are presented for confirmation
in (b). . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 59
4.13 (a) Ti-39Pt alloy heat treated at 1050 ◦C (HT2). Three phase are present in the
microstructure Ti4 Pt3 , Ti3 Pt, and Ti5 Pt3 . SADPs for Ti3 Pt, B=102, and Ti5 Pt3 ,
B=0001, are presented for confirmation in (b). In the case of the SADP from
Ti3 Pt, the 001 reflection is present due to double diffraction. (c) Ti-39Pt alloy heat
treated at 800 ◦C (HT3). Three phase are present in the microstructure Ti4 Pt3 ,
Ti3 Pt, and Ti5 Pt3 . SADPs for Ti3 Pt, B=102, and Ti5 Pt, B=2110, are presented
for confirmation in (d). In the case of the SADP from Ti3 Pt, the 001 reflection is
present due to double diffraction. . . . . . . . . . . . . . . . . . . . . . . . . . . . 60
ix
4.14 Ti-44Pt alloy heat treated at 1260 ◦C (HT1). The microstructure reveals the presence of two phases β-TiPt (it transforms to α-TiPt upon quenching) and Ti5 Pt3 .
The SADP from Ti5 Pt3 , B=2110, is presented for confirmation in (b). . . . . . . . 61
4.15 Ti-44Pt alloy heat treated at 1050 ◦C (HT2). (a) BEI reveals the presecence of four
phases: α-TiPt, Ti4 Pt3 , Ti5 Pt3 and a medium gray phase. In (b) TEM image of this
medium gray phase indicates that it is Ti4 Pt3 and the EDS spectra in (d) obtained
from the phases shown in Figure 4.15b revealed that this Ti4 Pt3 (II) adjacent to the
Ti5 Pt3 contains slightly more Ti compared to the matrix Ti4 Pt3 (I). . . . . . . . . . 62
4.16 Ti-44Pt alloy heat treated at 800 ◦C (HT3). (a) BEIs revealed the presence of
Ti5 Pt3 , α-TiPt, Ti4 Pt3 (I) and Ti4 Pt3 (II); the latter two observed also in the Ti-44Pt
alloy subjected to HT2. The BEI in (b) and the TEM images in (c-d) revealed the
presence of a fine lamellar structure of Ti4 Pt3 and α-TiPt that had transformed from
the prior β-TiPt lamellae observed at 1050 ◦C. . . . . . . . . . . . . . . . . . . . . 64
4.18 (a) BSE image indicating the location of the EDS line scan. (b) Results of the line
scan. Note that the platinum concentrations remain approximately constant along
the cross section of the dendrites. The mean Pt composition is 41.5 at.% Pt and the
standard deviation (SD) from the mean is 0.92. . . . . . . . . . . . . . . . . . . . 66
4.19 Modification to the Ti-Pt phase diagram appears as the dashed lines for the βsolidus and the β-solvus. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 67
4.21 Schematic sequence of transformations in the Ti-39Pt alloy during solidification.
The x-axis has its origin at the β-TiPt/Ti3 Pt interface and extends towards the core
of the dendrites. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 69
4.22 Schematic free energy curves for the transformations in the Ti-39Pt alloy during
solidification. The peritectoid transformation takes place at the peritectoid temperature (Tp ) involving Ti3 Pt of ∼ 30 at.% Pt and β-TiPt of ∼ 46 at.% Pt, reacting to
form Ti4 Pt3 of ∼ 42 at.% Pt. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 70
4.23 Schematic free energy curves for the transformations in the Ti-39Pt alloy during
solidification. At a temperature below the peritectoid temperature Tp , a β-TiPt
of 41.5 at.% Pt can transforms in a partitionless manner to Ti4 Pt3 of the same
composition. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 70
4.24 as-cast Ti-44Pt alloy. (a) BSE image indicating the location of the EDS line scan.
(b) Results of the line scan. Note that the platinum concentration decreases from
the core of the dendrite towards the periphery resulting in the different phases and
morphologies observed in the microstructure. . . . . . . . . . . . . . . . . . . . . 72
4.25 Schematic sequence of transformations in the Ti-44Pt alloy during solidification.
The x-axis has its origin at the β-TiPt/Ti3 Pt interface and extends towards the core
of the dendrites. Regions I, II and III correlate with the as-cast microstructure
presented in (d). . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 73
x
4.26 Samples of the Ti-39Pt alloy were heated to 1220 ◦C, 1240 ◦C and 1260 ◦C followed by an immediate WQ. (a) Heat treatments carried out to capture the transformation observed at around 1230 ◦C in the DTA scans. The microstructures in (b-c)
reveal the presence of a lamellar structure of Ti4 Pt3 +β-TiPt (which transforms to
α-TiPt upon quenching). . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 74
4.27 Schematic Gibbs free energy curves used to explain the sequence of transformations in the Ti-39Pt alloy during heat treatment. First, the peritectoid transformation takes place at T = Tp − ΔT (ΔT is the undercooling). . . . . . . . . . . . . . 75
4.28 Schematic Gibbs free energy curves used to explain the sequence of transformations in the Ti-39Pt alloy during heat treatment. Second the transformation of
β-TiPt to the lamellar Ti4 Pt3 +β-TiPt structure. . . . . . . . . . . . . . . . . . . . . 75
4.29 Ti-39Pt alloy subjected to heat treatment HT1 (Table 4.1) (a) the microstructure
reveals that some regions of the dendrites have transformed to 100% β-TiPt, while
others have transformed to the lamellar Ti4 Pt3 +β-TiPt structure (then β-TiPt transforms to α-TiPt upon quenching). (b) The peritectoid transformation can be observed along the interfaces Ti3 Pt/β-TiPt. . . . . . . . . . . . . . . . . . . . . . . . 76
4.30 (c) The heat treatment HT1 was repeated resulting in 100% α-TiPt dendrites (βTiPt→ α-TiPt). (d) After air cooled a sample heat treated HT1 the lamella Ti4 Pt3 +βTiPt structure is obtained in the prior β-phase. . . . . . . . . . . . . . . . . . . . . 77
4.31 SADPs reveal the OR between the Ti4 Pt3 and α-TiPt, which transforms from βTiPt upon cooling. (a) Transformation of β-TiPt to the lamellar Ti4 Pt3 +β-TiPt
structure. (b) Peritectoid reaction. Red fonts denote the reflections from the α-TiPt
phase, while the blue fonts denote the reflections from the Ti4 Pt3 phase. B= 0001
from Ti4 Pt3 (with pseudo-heaxoganl axes, see Chapter 3 for the charaterization of
this phase), and B= 110 from α-TiPt. . . . . . . . . . . . . . . . . . . . . . . . 78
4.32 Ti-39Pt sample heat treated at 1050 ◦C followed by an immediate water quenching.
The microstructure reveals that the α-TiPt layer, observed in the periphery of the
dendrites in the as-cast alloy, disappeared and possibly transformed to Ti4 Pt3 . . . . 79
4.33 DTA results on samples heated following the HT1 treatment. In (a) the initial
microstructure contains dendrites of 100% martensite while in (b) the initial microstructure contains the Ti4 Pt3 +α-TiPt lamellar structure within the dendrites. In
both experiments an exothermic peak is observed around 690 ◦C and is probably
associated with the transformation of the α-TiPt phase to Ti4 Pt3 . . . . . . . . . . . 80
4.34 (a) The microstructure of the sample heat treated at 670 ◦C followed by immediate
quench contains more Ti4 Pt3 within the dendrites. (b) The TEM image shows that
new Ti4 Pt3 (white box) is growing on the prior α-TiPt. . . . . . . . . . . . . . . . 81
4.35 (a) The microstructure of the sample heat treated at 1025 ◦C followed by immediate
quench reveals that the majority of the α-TiPt transformed into Ti4 Pt3 . (b) This
transformation results in Ti4 Pt3 of different variants within the dendrites. . . . . . . 81
4.36 Equilibrium condition for the formation of a Ti4 Pt3 of lower Pt content (X1 ) adjacent to Ti5 Pt3 in the Ti-44Pt alloys subjected to Ht2 and HT3 heat treatments. . . . 82
xi
4.37 Ti-44Pt alloy subjected to HT3 heat treatment. (a-b) A fine lamellar Ti4 Pt3 + α-TiPt
structure transformed in the prior β-TiPt lamellae observed in the microstructure
of the Ti-44Pt alloy subjected to HT2 heat treatment (Figure 4.15a). . . . . . . . . 83
4.38 Schematic showing the sequence of transformations in the Ti-31Pt alloy in the ascast condition (A), after HT1 (B), and after HT2 and HT3 (C). The sequence for
the Ti-35Pt is similar except that in the as-cast condition there is no primary phase.
84
4.39 Schematic showing the sequence of transformations in the Ti-39Pt alloy in the
as-cast condition (A), after HT1 (B), and after HT2 and HT3 (C). . . . . . . . . . . 85
4.40 Schematic showing the sequence of transformations in the Ti-39Pt alloy in the
as-cast condition (A), after HT1 (B), and after HT2 (C) and HT3 (D). . . . . . . . 86
A-1 Neutron diffraction pattern at three different temperatures: RT, 800 ◦C and 1100 ◦C in
˚ XRD pattern at RT of a Ti-42Pt as-cast powder sample. . . . . 108
the range 2.5-3.7 A.
xii
LIST OF TABLES
2.1
Summary of the of the known crystal structure, space group and lattice parameters
of the phases present in the composition range of interest of this work. . . . . . . . 11
3.1
Composition of the various phases observed in the microstructure of the Ti-42Pt ascast alloy heat treated at 1200 ◦C for 4 h determined using Wavelength Dispersive
Spectroscopy (WDS) in the EPMA. . . . . . . . . . . . . . . . . . . . . . . . . . 30
3.2
Platinum, titanium and oxygen content of the phases observed in the Ti-42Pt alloy
subjected to HT5 and HT6 heat treatments as determined by WDS in the EPMA. . 31
3.3
Lattice constants and space groups of the new phases observed in the as-cast and
heat treated alloy Ti-Pt alloys. . . . . . . . . . . . . . . . . . . . . . . . . . . . . 34
3.4
Crystallographic information of several Nowotny phases with the Mn5 Si3 prototype structure. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 34
4.1
Nomenclature and description of the heat treatments performed on the alloys. . . . 41
4.2
Chemical composition of the as-cast Ti-35Pt alloy measured at NASA Glenn Research Center using ICP-AES. . . . . . . . . . . . . . . . . . . . . . . . . . . . . 42
4.3
Platinum content of the phases observed in the as-cast Ti-31Pt alloy determined by
EDS in the FESEM. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 43
4.4
Pt content of the phases observed in the as-cast Ti-35Pt alloy as determined by
EDS in the FESEM. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 46
4.5
Pt content of the phases observed in the as-cast Ti-39Pt, 42Pt and 44P alloys determined by EDS in the FESEM. . . . . . . . . . . . . . . . . . . . . . . . . . . . 48
4.6
Platinum content of the phases observed in the Ti-31Pt alloys heat treated at 1260 ◦C,
1050 ◦C, and 800 ◦C. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 55
4.7
1050 ◦C, and 800 ◦C. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 57
4.8
1050 ◦C, and 800 ◦C. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 59
4.9
1050 ◦C, and 800 ◦C. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 63
4.10 Summary of microstructures observed in the Ti-31Pt alloy. . . . . . . . . . . . . . 87
xiii
4.13 Summary of microstructures observed in the Ti-39Pt alloy - Continued. . . . . . . 90
xiv
ACKNOWLEDGMENTS
I would like to thank my adviser, Professor Michael J. Kaufman, for his encouragement and
trust that he showed me throughout my academic career at Colorado School of Mines. His teaching
and advice have helped me to be a better professional every day. I will appreciate his mentoring
and support every day of my life.
My sincere thanks to Professor Stephen Liu, Professor Brian Gorman, Professor John Speer
and Professor Timothy Ohno for providing me with useful comments, and for being members of
my thesis committee.
I also acknowledge Dr. Ronald Noebe for providing me with samples, useful papers and data
necessary to develop this thesis.
I would also like to thank the members of my research group Grant Hudish, Mathew Kirsch
and Joseph Tsai because they all helped me when I needed it. I am also grateful to my colleagues
in the MME department.
This research and my graduate studies have been possible thanks to the support of a CONICYT
fellowship from the Government of Chile, a fellowship from the Universidad Técnica Federico
Santa Mar´ıa (Chile), the 2012 AIME Henry DeWitt Smith Scholarship and, the 2012 GSA Continuance Fellowship Award.
Thanks to Alejandro, my husband, who offered me assistance in numerous ways during my
work. Also, I would like to thank him for his patience and support every day that we have been
together, and above all for his love.
I want to express my gratitude to my dear parents, Cristóbal and Elsa, and my family for their
encouragement and support throughout my life.
xv
To Alejandro and
my beloved family
xvi
CHAPTER 1
INTRODUCTION
The search for materials capable of generating a mechanical response when subjected to a
thermal, electrical or magnetic field in applications that require lighter and stronger materials has
led to the development of “smart materials”. Shape Memory Alloys (SMAs) fall into the category
of smart materials and are capable of recovering their shape, after being deformed at low temperatures, when the temperature is increased. The leading SMAs on the market are those based on TiNi
alloys (Nitinols) which are used in dental devices, stents, bone plates and eyeglasses and, other applications where the application temperatures do not exceed 100 ◦C [4–6]. Given the potential
of SMAs at low temperatures, NASA and others have been developing High Temperature Shape
Memory Alloys (HTSMAs) for use in higher temperature applications such as sensors and actuators in jet engines, including, core exhaust chevrons, flow control devices, variable area/geometry
inlets, active clearance control devices, etc. [7].
Several alloying additions, namely Au, Pd, Pt, Hf and Zr, have been shown to increase the
martensitic transformation temperature responsible for the shape memory characteristic of TiNi
SMAs [7, 8]. Of these, Pt is known to have the greatest effect on the transformation temperatures.
Thus, several studies are ongoing to develop ternary Ti-Ni-Pt HTSMAs with an emphasis on the
composition range in the vicinity of that defined by Ti/(Ni+Pt)∼ 1 [9]. In the course of these
studies, it was discovered that there were previously unreported phases with compositions that
were Ti-rich relative to this 1:1 ratio where at least one of these phases is highly faulted with a
rather complex structure. TEM images of the quaternary as-cast alloy 51.8Ti-34.4Pt-6.5Ni-6.3Hf
(at.%) revealed the presence of a highly faulted phase (F) and martensite (M) in its microstructure
(Figures 1.1a and 1.1b). Therefore, it was decided that the Ti-rich side of the Ti-Pt binary system
should be reinvestigated to determine whether these were unique ternary phases or previously
unidentified phases in the binary Ti-Pt system.
1
M
F
(a)
(b)
Figure 1.1: Quaternary as-cast alloy 51.8Ti-34.4Pt-6.5Ni-6.3Hf (at.%) (a) BEI (b) TEM image
revealed the presence of a highly faulted phase (F) and martensite (M) in its microstructure.
Biggs and coworkers [1] carried out a detailed study of the Ti-Pt phase diagram in the composition range 30-60 at.% Pt (Figure 1.2). In this study, they reported the presence of a new phase in
alloys whose composition ranged between 30-47 at.% Pt and proposed that this new phase transforms via a peritectoid reaction (Ti3 Pt+β-TiPt → Ti4 Pt3 )
Figure 1.2: Ti-Pt phase diagram proposed by Biggs et al. (from [1]).
2
Among their observations, they showed:
i) The microstructure of an alloy of 30.4 at.% Pt revealed dendrites of Ti3 Pt, and Ti4 Pt3 in the
interdendritic regions.
ii) In alloys between 36.3 and 40.9 at.% Pt, they observed that the microstructure consisted
of dendrites of the Ti4 Pt3 phase and Ti3 Pt+Ti4 Pt3 eutectic in the interdendritic regions (Figure 1.3a).
iii) In addition, they found that the microstructure of the Ti-42.7Pt and Ti-43.1Pt (in at.%) as-cast
contained cored dendrites where the cores contained α-TiPt phase, in the form of irregular
laths, and the peripheries were mostly Ti4 Pt3 . Furthermore, the interdendritic regions were
comprised of Ti4 Pt3 and Ti3 Pt (Figure 1.3b).
(a)
(b)
Figure 1.3: SEM results reported by Biggs et al. [1]. (a) Backscattered Electron Image (BEI) of
Ti-37.6Pt as-cast alloy showing dendrites of Ti4 Pt3 with Ti3 Pt+Ti4 Pt3 eutectic in the interdendritic
regions. (b) BEI of Ti-42.7Pt as-cast alloy. The authors observed that the dendrites in (b) were
composed of cores of α-TiPt phase with Ti4 Pt3 at the peripheries.
iv) The authors suggested the existence of the Ti3 Pt(O) and Ti2 Pt(O) phases within the interdendritic regions of hypereutectic alloys that were heat treated at 1200 ◦C. Using Energy
3
Dispersive Spectroscopy (EDS) in the Scanning Electron Microscope (SEM), they suggested
that the Ti2 Pt(O) phase has a composition of 53Ti-29Pt-14O (at.%).
Since the peritectoid reaction must take place at the interfaces of the two parent phases (Ti3 Pt
and β-TiPt) and that the subsequent growth of this product phase (Ti4 Pt3 ) will be limited by diffusion of atoms from one parent phase to the other across the product phase, the rate of Ti4 Pt3 growth
via a peritectoid reaction must decrease with time/thickness. At lower temperatures, this rate may
become negligible [10]. However, the microstructure shown in Figure 1.3a is in contradiction with
this fact as the Ti4 Pt3 phase is not only observed at the dendrite peripheries (where the Ti3 Pt is
adjacent to the primary β-TiPt phase) but throughout the primary β-TiPt dendrites indicating that
the β-TiPt phase transforms directly to Ti4 Pt3 . Therefore, there is no conclusive evidence that
confirms that the peritectoid reaction takes place experimentally in alloys within this composition
range as proposed by Biggs et al. [1].
The morphology observed in the core of the dendrites of the Ti-42.7Pt as-cast alloy suggests
that there is some kind of partitioning of the prior β-TiPt phase taking place during solidification.
Since it is not possible to acquire reliable oxygen compositions using EDS, the composition of the
proposed Ti2 Pt phase (∼33 at.% Pt) in the interdendritic regions may indicate that the oxygen introduced by the titanium likely plays an important role in the phase observed and the microstructural
evolution of alloys from this portion of the diagram.
Biggs et al. also performed x-ray diffraction scan (XRD) on a specimen of the Ti-43Pt alloy
containing mostly Ti4 Pt3 . They concluded that the x-ray pattern belongs to a phase not observed
before, but, they did not determine the crystal structure of this unknown phase (Ti4 Pt3 ). Likewise, the authors did not perform any further characterization (e.g., TEM) that could allow the
unequivocal identification of the phases observed in the microstructures.
Therefore, the purpose of this research was i) to clarify the contradictions encountered in the
empirical observations of alloys in the composition range of interest (30-50at% Pt), ii) to confirm
the presence and nature of the peritectoid reaction, iii) to provide other corrections to the phase
4
diagram and clarify the role of interstitials on the various structures, and iv) to characterize the
crystal structures of the new phases present in the alloys of interest.
5
CHAPTER 2
BACKGROUND AND LITERATURE REVIEW
Shape Memory Alloys (SMAs) have the ability to recover (“remember”) their shape when the
temperature is increased. This property makes them useful for sensor and actuation applications in
industrial sectors such as aerospace, automotive, biomedical, and oil exploration [6]. The memory
effect in these alloys is related to the reversibility of the martensitic transformation proposed by
Kurdjumov and Khandros [11], who studied experimentally the reversible martensitic structure in
CuZn and CuAl alloys.
Near-equiatomic Ni-Ti alloys, known as Nitinols, are the SMAs most commonly utilized in
engineering applications and they continue to dominate the growing market due to their desirable
mechanical properties compared to other SMAs, good corrosion resistance, and biocompatibility.
Despite their advantages, the operating temperatures of the Nitinols fall below 100 ◦C making them
unsuitable for higher temperature applications [12]. In the early 1970s, High Temperature Shape
Memory Alloys (HTSMAs), namely titanium-platinum, titanium-palladium, and titanium-gold,
were developed and found to have martensitic transformation temperatures greater than 100 ◦C [6].
For example, Donkersloot and Van Vucht [8] studied Au-Ti, Pd-Ti, and Pt-Ti alloys near the
equiatomic composition. In each of these alloys the martensitic transformation temperature was
considerably above 500 ◦C (Figure 2.1) compared to the Ti-Ni equiatomic composition; in the
case of Ti-Pt, it was near 1070 ◦C. Considering these results, additions of Pt, Pd or Au to the Ti-Ni
alloys were found to increase considerably the transformation temperatures. The shape memory
effect is not limited to metallic systems; on the contrary, it has been observed in ceramics [13, 14]
and polymers [15, 16].
Further investigations on additions of alloying elements to the TiNi binary alloys revealed that
Pt additions have a remarkable effect on the martensitic transformation temperature, increasing the
6
temperature from approximately 50 ◦C to 1050 ◦C with 50 at.% Pt, see Figure 2.2 [7]. Therefore, Platinum has shown promising attributes towards increasing the transformation temperature
making these SMAs possibly suitable for high temperatures applications.
700
520
Temperature ◦ C
Ms
Mf
1080
Ms
Mf
600
500
500
480
400
460
300
440
200
420
Ms
Mf
1060
1040
1020
1000
980
960
100
400
40 42 44 46 48 50 52 54
40 42 44 46 48 50 52 54
at.% Au
at.% Pd
40 42 44 46 48 50 52 54
at.% Pt
Transformation Temperature ◦C
Figure 2.1: Alloying element effect on the martensitic start and finish transformation temperatures
for the binary alloys TiAu, TiPd, and TiPt [8].
Amount of Ternary Alloying Addition, at.% X
Figure 2.2: Transformation temperature versus percentage of alloying element added to Ti-50
at.%Ni alloy (from [7]).
SMAs possesses two remarkable characteristics, namely, the shape memory effect (SME) and
7
the superelastic effect (SE). In essence, both effects involve the transformation of a high temperature phase (austenite) to a low temperature phase (martensite) via a temperature-load cycle (SME)
or by applying a load in the austenitic region (SE). Both effects will be explained below. The transformation from austenite to martensite is diffusionless and occurs by shear distortion of the parent
lattice. Further, it is reversible, which makes the transformation useful for cyclic operations.
2.1
Ti-Pt phase diagram
Nishimura and Hiramatsu [2] carried out the first in-depth study of the Ti-Pt system, using
metallography, X-Ray Diffraction (XRD), melting point measurement, and differential thermal
analysis (DTA) and they constructed the Ti-Pt phase diagram shown in Figure 2.3a. Some important findings are summarized as follows:
• The presence of three intermetallic phases: Ti3 Pt, TiPt, and TiPt3 which melt congruently at
1370 ◦C, 1830 ◦C, and 1950 ◦C respectively.
• eutectic reaction: L → β-Ti + Ti3 Pt at 1310 ◦C and approximately 15 at.% Pt.
• eutectic reaction: L → β-TiPt + Ti3 Pt at 1320 ◦C and approximately 34 at.% Pt.
• eutectic reaction: L → β-TiPt + TiPt3 at 1780 ◦C and approximately 58 at.% Pt.
• eutectoid reaction: β-Ti→ α-Ti+Ti3 Pt at 840 ◦C and approximately 3 at.% Pt.
• peritectic reaction: L + TiPt3 → γ(Pt) at around 1800 ◦C and approximately 81 at.% Pt.
In 1965, Pietrowsky [17] identified the TiPt8 phase on the Pt-rich side of the phase diagram.
The author optimized an XRD pattern obtained using a Debye-Scherrer camera and CrKα radiation. Later, Junod et al. [18] studied experimentally the superconducting transition temperature and low temperature specific heat of binary and ternary compounds of titanium (Ti3 Pt, Ti3 Ir,
Ti75 Irx Pt25−x , and Ti3 Hg) with the A15-type structure (Pm3n) and reported that the Ti3 Pt phase
was not a stoichiometric phase as Nishimura and Hiramatsu had suggested; instead, they measured
a range of compositions between 22 ± 2 and 29 ± 2 at.% Pt at 500 ◦C.
8
Later, Murray [3] carried out thermodynamic calculations using the data compiled by previous
investigators [2, 8, 17, 18] and proposed the modified phase diagram shown in Figure 2.3b. This
diagram includes the composition range of the Ti3 Pt phase proposed by Junod et al. [18] as its
maximum composition range at 1310 ◦C. In addition, Murray included the two crystal structures
of the TiPt phase determined by Van Vucht and Donkersloot [8]: α-TiPt (Pmma with a = 0.455
nm, b = 0.273 nm and c = 0.479 nm), is stable from room temperature up to 1070 ◦C at 50
at.% Pt, and β-TiPt (Pm3m with a = 0.3192 nm), which is stable from 1070 ◦C up to the melting
temperature; these phases are related via the reversible martensitic transformation.
Recently, Biggs et al. [1] reported the presence of a new phase between the TiPt and Ti3 Pt
phases in contrast to the previous diagrams where no intermediate phases were present. They
carried out SEM, microprobe and chemical analyses, XRD, and DTA. The new phase observed
was reported to have the Ti4 Pt3 stoichiometry based on measured compositions in the range 41.743.4 at.% Pt. The results of XRD experiments were inconclusive and they were unable to determine
the structure of this phase. Their DTA results were used to establish the phase boundaries and to
conclude that the diagram published by Murray [3] was incorrect, and that the Ti4 Pt3 is a stable
phase which forms via a peritectoid reaction from the β-TiPt and Ti3 Pt phases at approximately
1205 ◦C, see Figure 2.3c.
More recently, Li et al. [19] developed a thermodynamic optimization of the Ti-Pt system
utilizing the Thermo-Calc software. The input to the optimization procedure considered thermodynamic data compiled from the literature [20] and the phase diagrams proposed by [1–3], and
the results were used to construct the phase diagram shown in Figure 2.3d. The crystal structure
of phases involved in the various transformations observed in the composition range of interest are
summarized in Table 2.1.
9
(a)
(b)
(c)
(d)
Figure 2.3: Phase Diagram of Ti-Pt system proposed by: (a) Nishimura and Hiramatsu in 1957 [2],
(b) Murray in 1982 [3], (c) Biggs et al. [1] and (d) Li et al. [19].
2.2
Shape Memory Effect (SME) and Superelastic Effect (SE)
In general, the SME and SE take advantage of the reversibility of the martensitic transfor-
mation (MT) which is a diffusionless transformation that involves the transformation of a high
temperature phase (austenite, A) to a low temperature phase (martensite, M). Below the characteristic transformation temperature, denoted Ms (s for starting), the free energy of the system
is lowered during the transformation A→M. This free-energy difference is the driving force for
the martensite reaction. The transformation upon cooling (forward transformation) begins at the
10
Table 2.1: Summary of the of the known crystal structure, space group and lattice parameters of
the phases present in the composition range of interest of this work.
Phase
Prototype
Space Group
Ti3 Pt
Cr3 Si
Pm3n
α-TiPt
AuCd
Pmma
β-TiPt
CsCl
Pm3m
Cell Parameters
lattice parameter (nm) angle
α = 90◦
β = 90◦
γ = 90◦
α = 90◦
β = 90◦
γ = 90◦
α = 90◦
β = 90◦
γ = 90◦
a = 0.5032
b = 0.5032
c = 0.5032
a = 0.455
b = 0.273
c = 0.479
a = 0.3192
b = 0.3192
c = 0.3192
Reference
[21]
[22]
[8]
martensite start temperature Ms and finishes at the martensite finish Mf when all the austenite phase
has transformed to martensite. Similarly, upon heating (reverse transformation), austenite begins
to form at the austenite start temperature As and the complete transformation ends at the austenite
finish temperature Af (Figure 2.4).
Austenite fraction
1
heating
cooling
0
Mf
As
Ms
Af
Temperature
Figure 2.4: Schematic of the martensitic transformation (MT) showing the characteristic temperatures involved during the transformation
Figure 2.5 shows the structural relationship between the austenite TiPt-β and the martensite
TiPt-α during the martensitic transformation which can be described as a shear shuffle of the
(110)B2 basal plane along the [110]B2 direction necessary to create the B19 structure [4].
One important characteristic of the martensitic transformation is the existence of a hystere-
11
Austenite B2
(β-TiPt)
Martensite B19
(α-TiPt)
(a)
(b)
Figure 2.5: Structural relationship between the parent B2 phase (TiPt-β) and the martensite B19
(TiPt-α). Martensitic transformation can be described as a shear shuffle of the (110)B2 basal plane
along the [110]B2 direction necessary to create the B19 structure [4].
sis during the thermal cycle as shown in Figure 2.4. The reverse transformation does not occur
in the same temperature interval as that in which the forward transformation occurs. Reversible
martensitic transformations (also known as thermoelastic transformations) are characterized by
small hystereses compared to a non-reversible martensitic transformation. For example, thermoelastic transformations in the Au-47.5 at.% Cd alloy exhibit a hysteresis of the order of 15 K, while
the non-thermoelastic transformation in the Fe-30 wt.% Ni alloy exhibits a hysteresis of the order
of 400 K [23, 24]. This large temperature difference is indicative of a large driving force needed
in this alloy to nucleate the transformation and the interface austenite/martensite is immobile once
the martensite grows to some size; therefore, making the reverse transformation irreversible. In
the former system, the interface is very mobile upon cooling and heating favoring the reversibility
of the martensitic transformation, so little driving force is required to accomplish the transformation [24].
In general, the MT in alloys occurs by shear distortion of the austenite lattice structure which
is achieved by the movement of the interface that separates the austenite phase from the martensite.
This shear takes place along a specific plane called the habit plane which is usually assumed to
be an undistorted plane. As the interface moves, atoms in the lattice structure of the austenite
phase are realigned into the lattice of the martensite phase. The phenomenological theory for
martensitic transformation states that to reduce the strains introduced to the martensite, an invariant
12
plane strain condition is required. However, a shear component is still not eliminated under this
condition and, therefore, a step of strain accommodation is necessary in SMAs. This is achieved
by the combination of favorable habit planes by a mechanism known as self-accommodation of
martensites. These habit planes are twin related to each other [6, 24]. This state of strain generated
in the lattice of the martensite could be either elastic or plastic or a combination of both, which is
an important consideration if the martensitic transformation is reversible or not.
If the strain generated in the martensite is elastic, the boundaries between the martensite and
the austenite phases are able to move both easily and reversibly. Thus, if martensite is formed
on cooling, then on reheating the specimen, the martensite reverts to the austenite. If recooled,
martensite reappears. This cycle can be repeated again and again, and no macroscopic shape
change is observed. On the contrary, if plastic deformation is introduced in the martensite, then the
boundaries between the martensite and the austenite become immobile by the dislocation structure
that results from the growth of the martensite. In this case, upon heating the austenite is forced to
nucleate inside the martensitic phase. This means that the martensitic transformation is irreversible.
2.2.1
Thermodynamics of the Thermoelastic Martensitic Transformation
The reversibility of the martensitic transformation (thermoelastic transformation) was predicted by Kurdjumov and Khandros [11]. Because martensitic transformations are not associated
with a compositional change, the free energy of the austenite and martensite phases can be represented schematically as is shown Figure 2.6, where GA is the free energy of the austenite phase,
GM is the free energy of the martensite phase, To is the equilibrium temperature between these two
phases, ΔGM →A is the driving force for the forward transformation, ΔGA→M is the driving force
for the reverse transformation, ΔTF is the undercooling required for the martensitic transformation
in the forward direction, while ΔTR is the superheating required for the martensitic transformation
in the reverse direction [24].
During a thermoelastic transformation, an equilibrium should exist between the chemical and
13
G
GA
GM
ΔGA→M
ΔGM→A
ΔTF
Ms
ΔTR
To
As
Temperature
Figure 2.6: Schematic representation of the free energy curves for the austenite and martensite
phases. ΔTF is the undercooling required for the martensitic transformation in the forward direction, while ΔTR is the superheating required for the martensitic transformation in the reverse
direction.
non-chemical forces. The equilibrium condition for the forward transformation from the austenite
to the martensite phase may be written [23]:
+ ΔGA→M
=0
ΔGA→M = −ΔGA→M
ch
nch
(2.1)
where ΔGA→M
is the chemical component of the Gibbs free energy and ΔGA→M
is the nonch
nch
chemical component. The non-chemical part can consist of a number of parts; the most important
and the work done against frictional forces, EfA→M
. Then,
are the elastic energy ΔGA→M
el
r
A−M
ΔGnch
= ΔGA−M
+ EfA−M
el
r
(2.2)
The elastic portion in Equation 2.2 is primarily due to the energy stored in the A/M interfaces
and to the elastic strains due to the creation of martensite, which has a slightly different volume
than the austenite. In turn, the frictional losses include the work done in moving the martensite
interfaces, the energy losses associated with the creation of internal defects induced by the transformation (such as stacking faults and twin boundaries), and the partial plastic accommodation
14
arising from volume and shape changes associated with the transformation. The most important
of these losses are those due to the motion of the interfaces. These frictional losses result in the
thermal hysteresis loss.
2.2.2
Shape Memory Effect
According to Kumar and Lagoudas [6], when the austenite phase is cooled below the Mf
temperature, the martensitic transformation results in the formation of several martensitic variants which arrange in such a way that the macroscopic shape change is negligible. This form of
martensite is known as twinned martensite. A schematic of this transformation path is shown in
Figure 2.7a. If a load above the σs (detwinning start stress) is applied to the twinned martensite
at temperatures below Mf the martensite can be “detwinned”, and therefore, several variants can
reorient into a preferred variant resulting in a detwinned martensite. In this case, a macroscopic
shape change occurs and the deformed configuration is retained after the load is released. When
the detwinned martensite is heated to a temperature above Af the crystal structure transforms back
to austenite, and the original shape of the SMA is regained. The cycle can be repeated obtaining
again twinned martensite below the Ms temperature. This second path is called the Shape Memory
Effect (SME) (or one-way SME) and is presented in Figure 2.7b.
2.2.3
Superelastic Effect (SE)
The SE is associated with the stress-induced transformations of austenite to martensite. It occurs when the austenite phase is subjected to a mechanical load at temperatures slightly above the
Af . By applying a stress above the σf (detwinning finish stress), the austenite transforms to a fully
detwinned martensite. If the temperature of the material is above Af , a complete shape recovery
is achieved upon unloading the specimen when martensite returns to austenite. A common pseudoelastic path is schematized in Figure 2.8a in which an isothermal superelastic test is performed.
This path can be explained with the help of a stress-strain curve (Figure 2.8b): i) A→B The load is
15
(a)
(b)
Figure 2.7: (a) Schematic of the twinned martensite which transforms upon cooling in the absence of an applied load, while (b) shows a schematic of the detwinned martensite which forms at
low temperatures when a sufficient high load is applied to reorient the martensitic variants into a
preferred orientation (from [6]).
applied to the austenite which, in turn, undergoes elastic deformation. At σM s the austenite phase
starts to transform to martensite. ii) From B→C the martensitic transformation proceeds until the
stress σMf is achieved. At this point the MT ends. iii) C→D Subsequent increase in the stress
does not cause further transformation, only elastic deformation of the martensite. iv) D→E The
martensite begins to unload elastically. At the stress σAs the reverse transformation takes place and
16
austenite begins to form. v) E→F The reverse transformation continues until at σAf the martensite
phase has fully transformed to austenite. vi) F→A The specimen unloads elastically until the stress
is zero.
(a)
(b)
Figure 2.8: The austenite phase existing at a temperature above Af is subjected to a stress above
σMf resulting in fully detwinned martensite. When the specimen is unloaded the austenite phase
is fully recovered. This is known as superelastic effect (from [6]).
17
2.3
Organization of the thesis
The remainder of this dissertation is organized into two major chapters. Chapter 3 will treat
the characterization of the new phases observed in the microstructure of the as-cast and heat treated
alloys presented in Chapter 4. Selected Area Diffraction (SADP) and Convergent Beam Electron
(CBED) patterns were obtained using transmission electron microscopes located both at the Colorado School of Mines and at the NASA Glenn Research Center
Chapter 4 will treat the results related with the findings encountered in as-cast and heat treated
alloys whose composition ranges between 30-50 at.% Pt and their relationship with the phase
diagram proposed by Biggs et al. [1]. The discussion will be centered on the possible path of phase
transformations and a possible modification of the phase diagram. Finally, an overall conclusion
and discussion will be provided in Chapter 5.
18
CHAPTER 3
CHARACTERIZATION OF PHASES
The microstructures from the as-cast and heat treated alloys revealed the presence of phases
that have not been characterized previously such as Ti4 Pt3 , and other phases that are new in the
Ti-Pt phase diagram such as Ti5 Pt3 , a globular phase observed in the low platinum alloys (Ti/Pt
ratio∼4), both of which seem to be stabilized by interstitials as it will be shown below. The purpose
of this chapter is to report the structures and compositions of these phases. In addition, another
phase was also identified and appears to also be due to contamination.
3.1
Experimental Methodology
Six alloys, namely Ti-34Pt, 39Pt, 42Pt, 44Pt and 50Pt (at.% Pt) were prepared at the NASA
Glenn Research Center by non-consumable arc melting high purity starting components Ti (99.995
wt.%) and Pt (99.995 wt.%) together in the appropriate proportions. Since platinum (21.45 g/cm3 )
is almost five times denser than titanium (4.5 g/cm3 ) and has a slightly higher melting temperature,
the buttons were flipped and remelted several times in order to insure complete melting and mixing.
In spite of these efforts to insure mixing, problems were occasionally encountered in the form of
unmelted Pt and this was taken into account in the different analyses.
While various characterization methods were employed, the structural analysis was done
mostly using analytical and high resolution transmission electron microscopy. A Phillips 400T
and a Phillips CM12 transmission electron microscopes operated at 120 kV accelerating voltage,
and a Phillips CM200 operated at 200 kV accelerating voltage were employed for these analyses.
SEM/EDS analyses were also used to characterize these alloys in order to obtain an indication of
the volume fraction, solidification structure and compositions of the various phases. Finally, electron microprobe analysis of the Ti-42Pt alloy was performed using a JEOL JXA 8200 Super Probe
19
in order to determine the compositions of the various phases more accurately. Simulations of the
diffraction patterns was also performed using the commercial JEMS software [25].
3.2
Results
3.2.1
Ti-34Pt Alloy
According to the Ti-Pt phase diagrams proposed by Murray [3] and Biggs et al. [1], the Ti34Pt alloy falls at approximately the eutectic between Ti3 Pt and TiPt (Figures 2.3b and 2.3c). The
as-cast microstructure of this alloy (Figure 3.1a) is clearly inconsistent with the published phase
diagram since, instead of the expected eutectic-like structure, there is a considerable volume fraction of primary dendrites of a low atomic number phase with interdendritic regions that might
be eutectic in nature although they do not appear to be classical coupled growth regions. EDS
analysis of the primary dendrites in the SEM revealed that they are Ti3 Pt while the higher Z interdendritic regions are comprised of α-TiPt and Ti4 Pt3 (Figure 3.1b) which is consistent with the
phases observed in the interdendritic regions of the Ti-35Pt and Ti-39Pt as-cast alloys (Figures 4.5b
and 4.6b).
TEM analysis revealed that the interdendritic regions were complex and consisted of multiple
phases (Figure 3.1c). Electron diffraction analysis revealed that, in addition to Ti3 Pt, there were
four different phases in these regions. These include the α-TiPt (martensite) and three other “new”
phases. Two of these three phases were found to have Pt contents intermediate to both TiPt and
Ti3 Pt while the third (small globular phase) had a Ti/Pt ratio of ∼4. In addition, this analysis
confirmed the presence of the cubic (a = 0.503 nm) Ti3 Pt phase (Figure 3.2). The characterization
of these new phases will be summarized below followed by a discussion of the microstructural
evolution and the relationship to the published and recently-modified phase diagrams.
20
Ti3 Pt
Ti3 Pt
α-TiPt+
Ti4 Pt3
50 μm
(a)
(b)
α-TiPt
Ti5 Pt3
Ti4 Pt3
globular
phase
200 nm
(c)
Figure 3.1: Low magnification (a) and high magnification (b) BEIs of the Ti-34Pt as-cast alloy
showing the primary dendrites of Ti3 Pt (dark phase) and multi-phase interdendritic eutectic comprised of Ti4 Pt3 and α-TiPt (which transforms from β-TiPt upon cooling). (c) Bright Field TEM
(BFTEM) of the interdendritic regions of the Ti-34Pt as-cast alloy. As can be seen there are multiple phases in these regions including three that are actually new phases, namely Ti5 Pt3 , Ti4 Pt3 and
the globular phase.
3.2.1.1
Faulted Phase: Ti4 Pt3 This highly faulted phase is observed in both the
hypoeutectic alloys (in the eutectic regions) and hypereutectic alloys (in both the eutectic and
21
m
m
m
m
m
(a)
(b)
100
010
B=001
(c)
Figure 3.2: Convergent Beam Patterns (CBPs) and SADPs of the Ti3 Pt phase in Figure 3.1a: (a)
011 CBP and (b) 111 CBP and (c) 001 SADP. These patterns are consistent with the cubic
structure of Ti3 Pt (a = 0.503 nm).
dendrites). The composition of this phase was analyzed using EDS in the TEM and was found to be
approximately 36 at.% Pt although, as will be shown below, Wavelength Dispersive Spectroscopy
(WDS) measurements indicated a higher Pt concentration (∼ 41 at.% Pt). Thus, this appears to
be consistent with the Ti4 Pt3 phase (41.7 – 43.4 at.% Pt) that was reported by Biggs et al. [1].
The analysis of this phase was rather cumbersome because of the faulting and the large lattice
parameters.
In such heavily-faulted structures, the symmetry is difficult to resolve in convergent beam
patterns and, therefore, one has to rely on careful tilting experiments and/or finding regions that
22
contain fewer or no faults. Even so, by using a combination of selected area electron diffraction,
convergent beam electron diffraction and careful tilting experiments, it was determined that the
heavily-faulted phase had a complex structure with considerable streaking in the resulting selected
area diffraction patterns (SADPs). The latter indicates the likely presence of planar faults (e.g.,
twins, antiphase boundaries, or polytype boundaries). Another example of this type of highly
faulted phase is encountered in the Ni-Co binary system with a cubic structure [26]. SADPs from
some of the major zones are shown in Figure 3.3 and it is apparent that the spot spacings are
very close together indicative of the large d-spacings/lattice constants mentioned above. As can
be seen in this figure, the closely-spaced systematic row that includes the transmitted beam is not
as heavily streaked as the adjacent parallel rows. This commonly occurs in materials that contain
a high density of closely-spaced twins where the unstreaked row corresponds to the twin plane
common to both matrix and twin or to the plane that is common to the two variants on either side
of the boundary. The spacings of these closely-spaced spots correspond to a large d-spacing of
∼ 2.36 nm.
When the sample was tilted along the Kikuchi band orthogonal to this row of spots, the major
zones were found to repeat approximately every 60◦ , which is typical of hexagonal or trigonal
structures with 6-fold or 3-fold symmetry, respectively. Thus, the structure was tentatively indexed
with hexagonal indices of a ∼ 0.8 nm and c ∼ 2.36 nm. The spacings of the rows (rather than the
spots within these rows) orthogonal to the “0001” systematic row are consistent with this indexing
scheme as indicated in the Figure 3.3.
In an effort to determine the point and space groups of this faulted phase, a considerable
amount of material was examined in an attempt to find regions free of faults. Fortunately, such regions were observed occasionally (Figures 3.4a and 3.4b) and, SADPs obtained from these regions
(Figures 3.5a and 3.5b) were found to contain readily resolvable diffraction spots that were not
obscured by the streaking. Close examination of the 1120 and 1100 SADPs revealed that there
are no systematic rows of diffraction spots that are orthogonal to the 000l row as would be the case
if the structure was hexagonal or trigonal (with the 3m, 3m or 32 point groups). If the structure was
23
Indices
d-spacing (nm)
(1100)
0.68
(1120)
0.40
(1230)
0.26
0.68 nm
B= 1120
0.26 nm
0.40 nm
B= 1100
B= 1340
Figure 3.3: SADPs from the faulted Ti4 Pt3 phase in Figure 3.1c. As discussed in the text, the
spacing of the spots in the unstreaked row and the spacings orthogonal to this row are readily
indexed using hexagonal indices with a ∼0.8 nm and c ∼ 2.36 nm (included in the table).
trigonal, e.g., 3m, then the 1120 patterns would not contain orthogonal rows whereas the 1100
patterns would.
Furthermore, the 1100 Convergent Beam Pattern (CBP) (Figure 3.6a) has no horizontal or
vertical mirror. Therefore, it is concluded that the structure of the Ti4 Pt3 phase, at least at lower
temperatures, does have a low symmetry and that the faulting in this pseudo-hexagonal phase is
related to the transformation from the higher symmetry β-TiPt phase where the faulting appears to
be related to the twin-related variants of the product phase. This assumption is based on the lack
of streaking along the 000l systematic rows in the various [uvi0] SADPs.
24
Faulted Ti4 Pt3
Faulted region
α-TiPt
Unfaulted region
Unfaulted Ti4 Pt3
100nm
5 nm
(a)
(b)
Figure 3.4: BFTEM (a) and high resolution electron micrograph (HREM) (b) of a region of the
faulted phase in the Ti-34Pt as-cast alloy that appears to contain no faults. This was rare as,
typically, this phase is heavily faulted throughout.
(a)
(b)
Figure 3.5: SADP of the [1120] (a) and [1100] (b) zones (referred to the hexagonal indices) of an
unfaulted region of the faulted phase revealing the lack of a mirror plane in the zero order Laue
zone. Thus, the symmetry of this pattern is either 2 or 1.
The lattice parameters are unknown but must be related to the pseudo-hexagonal parameters
where a ∼ b ∼0.80 nm and c ∼2.36 nm. The [0001] SADP from this phase (Figure 3.6b) appears
to be 3-fold or 6-fold with no streaking in the SADPs consistent with the faulting on the basal
25
planes lying orthogonal to the electron beam. The volume of this cell is ∼ 1.3 nm3 which, when
compared to that of the B2 structure (∼ 0.27 nm3 with 2 atoms per cell), indicates that this structure
probably has on the order of 96 atoms per unit cell. Clearly, further work is required to determine
if this assumption is correct and to identify the actual structure of this phase.
Basal Kikuchi band
(a)
(b)
Figure 3.6: (a) 1100 CBP of an unfaulted region of the “faulted” pseudo-hexagonal phase Ti4 Pt3 .
While it appears that there might be two orthogonal mirrors, the lack of a 0001 mirror in the 1120
pattern (Figure 3.5a) and the observed Kikuchi bands at high angles (marked) that are not mirrored
to the bottom indicate that the whole pattern symmetry is 1-fold. (b) SADP of the faulted phase
taken near B= 0001. As can be seen, the spot spacings and the angles between the systematic
rows are very close to those expected for a 3-fold or 6-fold axis.
3.2.1.2
Ti5 Pt3 Phase This unknown phase was observed in the heat treated alloys
and, also, in the Ti-42Pt and Ti-44Pt as-cast alloys. It contains a slightly lower platinum content
than the Ti4 Pt3 phase and can be seen as the medium gray phase in the BEIs presented in Sections 4.2 and 4.3. Convergent beam and selected area electron diffraction were used to analyze
this phase and the results are shown in Figures 3.7a to 3.7c. First of all, the CBPs in Figure 3.7a
have 6mm projection and whole pattern symmetries and, therefore, a diffraction group of either
6mm or 6mm1R . Likewise, the CBPs in Figures 3.7b and 3.7c have 2mm projection and whole
26
pattern symmetries and a corresponding diffraction group of 2mm or 2mm1R . Using the tables in
Buxton et al. [27], these two zones can be used to determine that 6/mmm is the only consistent
point group for this structure.
(a)
(b)
Figure 3.7: CBPs from the Ti5 Pt3 phase in the Ti-Pt alloys along (a) 0001, (b) 1120 and (c)
1100 directions. Note that the patterns in (a) are 6mm while those in (b) and (c) have 2mm
symmetry in the whole pattern consistent with the 6/mmm point group. Also, note the dynamic
absences (arrowheads) in (c).
27
0002
1120
(c)
Figure 3.7: Continued
The space group of this hexagonal phase was readily determined by noting the various dynamic absences and kinematic extinctions. For example, the [1100] CBP (Figure 3.7c) contains
dynamic absences in the 000l (l = 2n + 1) reflections. This is consistent with either a two-fold
screw axis (or equivalent 6-fold screw axis) along [0001] and/or a {1120} glide plane. Further, the
extinctions in Figure 3.8c are clearly of the type 000l, l = 2n + 1 and hh0l, l = 2n + 1. This is
different than the standard extinction conditions for typical hexagonal metals and alloys (e.g. Ti)
which have the P63 /mmc space group with extinctions of the type hh2hl, l = 2n + 1. Thus, considering the information revealed by the CBPs (Figures 3.7a to 3.7c) and the SADPs (Figures 3.8a
to 3.8c) the space group of the current phase is P63 /mcm. Based on this space group, the lattice
constants are readily determined to be a ∼ 0.8 nm and c ∼ 0.5 nm. Using EDS data, the composition of this phase was found to be approximately 37 at.% Pt, intermediate to the Ti3 Pt and Ti4 Pt3
phases.
A sample of the Ti-42Pt alloy was heat treated to 1200 ◦C for 4 hours in order to coarsen and
equilibrate the structure. The resulting microstructure (Figure 3.9a) appeared considerably differ-
28
1100
00
0002
1120
11
0002
(a)
(b)
(c)
Figure 3.8: SADPs from the Ti5 Pt3 phase in the Ti-Pt alloys along (a) 0001, (b) 1120 and (c)
1100 directions used to determine the lattice parameters and extinctions of this structure.
ent than the as-cast material (Figures 4.7a to 4.7c) and contained an intermediate Z phase in the
interdendritic regions, Ti5 Pt3 , and small particles of a low Z number phase. A high magnification
BEI of a sample of this alloy heat treated at 1260 ◦C for 2 hours followed by WQ is presented in
Figure 3.9b to confirm the presence of the small particles of low Z number. In an effort to obtain
a more quantitative analysis of the compositions of these new phases, this sample was examined
by electron microprobe analysis. The results are shown in Table 3.1. As can be seen, the matrix
composition (point 3) is from the Ti4 Pt3 phase and is measured to be approximately 59Ti-41Pt
which is richer in Pt than was measured using EDS above. Likewise, the high Z phase (point 5)
appears to be slightly enriched in Pt, consistent with it being the martensitic phase, α-TiPt, with
some percentage of the faulted phase. The interdendritic regions contain 3 phases all of which
are depleted in Pt relative to the matrix phases. The highest Z phase in this region is the Ti5 Pt3 ,
hexagonal P63 /mcm, phase with a composition Ti-36.2Pt. The intermediate Z phase corresponds
to Ti3 Pt phase with a composition Ti-29.9Pt.
Electron Probe Micro-Analysis (EPMA) results for the low Z phase in the interdendritic regions was measured to be Ti-21.8Pt-3.2C-15.4N. This phase corresponds to the Ti4 Pt phase observed also in the as-cast alloys. The compositional results indicate that there is a significant
interstitial content in this phase, which is in agreement with the EDS spectrum collected from this
29
phase using the EDS in the TEM (Figure 4.4c).
Ti3 Pt
Ti5 Pt3
α-TiPt
Ti4 Pt
(a)
(b)
Figure 3.9: Ti-42Pt alloy (a) heat treated at 1200 ◦C for 4 h followed by WQ. The microstrcuture
reveals the presence of five phases, namely Ti3 Pt, Ti4 Pt3 , α-TiPt, Ti5 Pt3 and Ti4 Pt. WDS measurements reveal that the small particles of low Z number contain a significant amount of interstitials
(C and N) and that corresponds to the “Ti4 Pt” phase. This result suggests that this phase is stabilized by interstitials during solidification and heat treatment. (b) Heat treated at 1260 ◦C for
2 h+WQ containing primary dendrites of β-TiPt (which transformed to α-TiPt upon quenching),
Ti3 Pt, Ti5 Pt3 and Ti4 Pt.
Table 3.1: Composition of the various phases observed in the microstructure of the Ti-42Pt as-cast
alloy heat treated at 1200 ◦C for 4 h determined using Wavelength Dispersive Spectroscopy (WDS)
in the EPMA.
Location
Contrast
Ti
Pt
C
N
Ti/Pt
Phase
1
light gray
63.8
36.2
-
-
1.8
Ti5 Pt3
2
dark gray
70.1
29.9
-
-
2.3
Ti3 Pt
3
matrix
58.8
41.2
-
-
1.4
Ti4 Pt3
4
black
59.6
21.8
3.2
15.4
2.7
“Ti4 Pt”
5
white
56.5
43.5
-
-
1.3
α-TiPt
Additional WDS measurements were carried out on one Ti-42Pt alloy heat treated at 1300 ◦C for
5 h followed by WQ (HT5), and another sample of the same alloy heat treated at 1100 ◦C for 168
h followed by a second heating at 1300 ◦C for 240 h (HT6) to determine the oxygen content of
30
the different phases present in these heat treated microstructures. The microstructures reveal the
presence of Ti3 Pt, Ti4 Pt3 , α-TiPt, Ti4 Pt and Ti5 Pt3 . The results are summarized in Table 3.2.
Table 3.2: Platinum, titanium and oxygen content of the phases observed in the Ti-42Pt alloy
subjected to HT5 and HT6 heat treatments as determined by WDS in the EPMA.
Phase
3.2.1.3
at.%
Pt
Ti
O
Ti3 Pt
29.2
70.6
0.2
Ti5 Pt3
34.9
60.8
4.6
Ti4 Pt3
41.1
58.2
0.7
α-TiPt
46.6
53.4
0.0
Globular Ti4 Pt(C,N) Phase This phase was observed in the Ti-31Pt and Ti-
34Pt as-cast alloys (Figures 3.1c and 4.4b), and in the Ti-42Pt heat treated alloy. SAD and CBED
analysis of this small globular phase revealed that its structure was inconsistent with anything in
the Ti-Pt diagram. The CBP in Figure 3.10a has 4mm whole pattern symmetry and either the 4mm
or 4mm1R diffraction group while that in Figure 3.10b has 3m whole pattern symmetry and the
6R mmR diffraction group. The CBP in Figure 3.10c has 2mm whole pattern symmetry and either
the 2mm or 2mm1R diffraction group. Using this information, it is possible to determine the point
group of this phase as m3m.
The space group of this phase was determined by noting that the structure was fcc-like except
that all the odd reflections (e.g., 111) were considerably weaker than the all even reflections (e.g.,
200) (e.g., see Figure 3.10d). This is common for compounds with the NaCl (B1) structure where
the structure factors for odd reflections may be considerably lower than for even reflections (for
NaCl, F = 4(fNa + fCl ) for even reflections and F = 4(fNa − fCl ) for odd reflections). Thus, the
space group of these small globular particles is Fm3m and the structure probably has a B1 structure similar to NaCl. Based on this analysis, the lattice parameter of this phase was measured and
found to be approximately 0.8 nm. It should be noted that the Ti(C,N) phase frequently observed
31
(a)
(b)
200
111
022
(c)
(d)
Figure 3.10: (a) 001, (b) 111 and (c) 011 CBPs from the globular Ti4 Pt phase. The symmetries of these patterns are consistent with an m3m point group and the observed forbidden reflections with an Fm3m space group. The weak odd reflections in the 011 SADP in (d) indicate a B1
(NaCl) ordering.
in Ti alloys, or in Ni-base superalloys that contain Ti, has the NaCl structure but the odd reflections are relatively strong due to the low atomic scattering factor of the interstitial elements for
electrons which do not greatly impact the relative intensities of the odd and even reflections (i.e.,
32
4(fTi +fC )∼4(fTi -fC )). In addition, the lattice parameter of this phase is normally slightly larger that
0.4 nm as opposed to 0.8 nm which indicates that the small globular particles in the interdendritic
regions are a different phase.
EDS analysis of this phase revealed that the Ti/Pt ratio was ∼4 which is higher than that of
any of the other phases. Since the primary phase in this alloy is the Ti3 Pt phase, it is concluded
that this phase is an interstitial phase that presumably forms from the last liquid which rejects the
interstitial elements from the other phases. The size and location of these small globular particles
is consistent with this scenario.
3.2.2
Discussion
The presence of an intermetallic phase corresponding to a Ti4 Pt3 type of stoichiometry between Ti3 Pt and TiPt confirms the report of an intermediate phase by Biggs et al. [1]. In addition
a second intermetallic phase, Ti5 Pt3 , and an interstitial phase Ti4 Pt(C,N) have also been observed
in this portion of the phase diagram and greater detail concerning the crystal structures of all three
phases have been provided above. These phases and their structures are summarized in Table 3.3.
As mentioned previously, neither of the two binary intermetallic phases exists in the Ti-Pt phase
diagram published by Murray [3].
The Ti5 Pt3 phase was present in some of the as-cast alloys (42Pt and 44Pt) and in the heat
treated alloys. This suggests that its formation requires the presence of interstitial atoms, especially oxygen, to allow its stabilization in the microstructure. This is achieved when the remaining
liquid gets enriched in impurity atoms during solidification and, when the diffusion of interstitial
atoms increases with increasing temperature during heat treatment. In general, the microstructures
revealed that the Ti5 Pt3 phase nucleates and grows mostly on Ti3 Pt/Ti4 Pt3 interfaces.
Biswas and Schubert [28] reported what are likely isostructural phases with the Mn5 Si3 prototype structure (Table 3.4), namely Zr5 Pt3 , Zr5 Ir3 and Hf5 Ir3 , with lattice parameters a ∼ 0.80
nm and c ∼ 0.54 nm. These authors were the first to report compounds showing the Mn5 Si3 -type
33
Table 3.3: Lattice constants and space groups of the new phases observed in the as-cast and heat
treated alloy Ti-Pt alloys.
Phase
Space
Group
Lattice parameters
(prototype)
(nm)
Composition
Comments
Ti4 Pt(C,N)Fm3m
(NaCl)
a ∼ 0.8
Ti-22Pt-3C-15N
New interstitial phase with
Fm3m structure
Ti5 Pt3
P63 /mcm
(Mn5 Si3 )
a ∼ 0.8, c ∼ 0.5
Ti-36Pt
Isostructural with Zr5 Pt3
phase
Ti4 Pt3
Unknown
a ∼ 0.8, c ∼ 2.36
(pseudo-hexagonal
indices)
Ti-41Pt
Faulted phase with
pseudo-hexagonal
structure
structure in which both components are transition elements. Later, Cenzual and Parthé [29] found
that Zr5 Ir3 and Hf5 Ir3 posses a superstructure of the hexagonal Mn5 Si3 phase where c = 3cMn5 Si3 .
They refined the structure and concluded that these compounds belong to the P61 22 (hP48) space
group that contain 48 atoms. The atoms positions for these phases with P61 22 (hP48) space group
are summarized in Table 3.4.
Table 3.4: Crystallographic information of several Nowotny phases with the Mn5 Si3 prototype
structure.
Phase
Prototype Space Group
Mn5 Si3 Mn5 Si3
Zr5 Ir3
(Pearson Sym.)
a,b,c
˚
A
atom positions
atom loc.
P63 /mcm
a=b=6.912
(hP16)
c=4.812
x
Ref.
y
z
occ.
Mn1 4d 0.3333
0.6667
0.0000
1.00
Mn2 6g 0.2360
0.0000
0.2500
1.00
Si1
6g 0.5595
0.0000
0.2500
1.00
P61 22
a=b=7.9306
Zr1 12c 0.2384
0.0062
0.3396
1.00
(hP48)
c=17.01
Zr2
6a 0.2480
0.0000
0.0000
1.00
Zr3 6b 0.3710
0.7420
0.2500
1.00
Zr4 6b 0.6557
0.3114
0.2500
1.00
Ir1
12c 0.4137
0.0817
0.1444
1.00
Ir2
6a 0.6111
0.0000
0.0000
1.00
34
[30]
[29]
In addition, the same Mn5 Si3 -type structure have been observed in different alloy systems
containing rare-earth elements [31–35]. Phases with this type structure are known as Nowotny
phases because extensive studies on this phase were carried out by Nowotny and coworkers [36–
41].
More recently, Leonard et al. [42] reported an ω phase with P63 /mcm space group in ternary
Nb-Ti-Al alloys with somewhat similar lattice constants (a ∼ 0.796 nm and c ∼ 0.557 nm). While
it is not clear whether this is an isostructural phase, it is tempting to speculate that it might be
related to the Ti5 Pt3 phase. Having said that, there is no indication in the present study that Ti5 Pt3
phase forms from the cubic B2 β-TiPt phase as is clearly the case in the Nb-Ti-Al alloy.
Several compounds with the stoichiometry A5 B3 and Mn5 Si3 -type structure can be obtained
in its pure form or stabilized by interstitials atoms, namely B, C, N, O and other elements, without
changing the basic structure and space group. Some examples of interstitial stabilized compounds
include Zr5 Sb3 X (X=C, O, Zn, Si), (Ca,Ba)5 Sb3 Cl, Zr5 Sn3 X (X=C, O, Ge, Ga), Zr5 Sn3 X (X=B, C,
N, O, Cu), Zr5 Pb3 Zn, La5 Ge3 X (X=N, O, Cr), Ca5 Pb3 X (X=Mn, Fe) and some La5 Ge3 Z [43]. The
fact that the Mn5 Si3 -type structures can act as a host for interstitials elements is interesting in light
of the microstructural observations made in Chapter 4, and allows one to conclude that Ti5 Pt3 O
may be classified as a “stuffed” phase with the Mn5 Si3 -type structure.
The crystal structure characterization carried out in the previous sections revealed that the
Ti4 Pt3 and Ti5 Pt3 phases are structurally related. The microstructures of the as-cast and heat treated
alloys confirmed that Ti5 Pt3 transforms from Ti4 Pt3 after sufficient oxygen content is present to allow its stabilization. The results presented in Chapter 4 revealed that Ti4 Pt3 is able to transform
directly from β-TiPt either via partitionless transformation during solidification (β-TiPt → Ti4 Pt3 )
or with partition of β-TiPt → β-TiPt+Ti4 Pt3 at approximately 1230 ◦C, depending on the composition of the β-TiPt phase. However, as noted above, the SADPs from the faulted Ti4 Pt3 phase are
such that they appear to have originated from a hexagonal structure at high temperatures. Additional studies are necessary to confirm if this is possible.
35
These two phases share the same a ∼ 0.8 nm lattice parameter. Closer examination of the
patterns in Figure 3.3 reveals a series of strong reflections that do appear to lie in orthogonal rows
as expected for a hexagonal structure (these are highlighted with a red box). Thus, an attempt was
made to overlay the basal and non-basal patterns from the common zones and the results shown in
Figures 3.11a to 3.11c are supportive of this speculation, i.e., the reflections in the 0001, 1123
and 1122 SADPs of the Ti5 Pt3 phase do indeed overlap with the strong reflections in the Ti4 Pt3
0
0
110
0 11
11 0
01
1 1
1
phase.
B = 1123
Ti5 Pt3
1 100
(b)
023
1123 2
(a)
Ti5 Pt3 +Ti4 Pt3
B = 1122
Ti5 Pt3
Ti5 Pt3 +Ti4 Pt3
(c)
Figure 3.11: (a) B = 0001 of both Ti5 Pt3 and Ti4 Pt3 phases, (d) B = 1123 of Ti5 Pt3 (left side)
and both phases (right side) and (e) B = 1122 of Ti5 Pt3 (left side) and both phases (right side).
36
Several A4 B3 -type structures are reported in the literature. Wilson et al. [44] refined a powder
diffraction obtained from Zr4 Al3 using CuKα radiation; they determined that this phase have a
˚ and c = 5.390 A.
˚ It transforms
hexagonal crystal structure with P6 space group and a = 5.433 A
via a peritectoid transformation [45]. Palenzona and Iandelli [46] examined several rare-earth
compounds of the type M3 Pd4 (M=Y, Th, Hf and U) and determined that they crystallized with the
˚ and a ∼ 5.8 A.
˚ FurPu3 Pd4 -type structure with R3 space group with lattice parameters a ∼ 13 A
thermore, Cenzual et al. [47] corrected the crystal structure of several compounds with monoclinic
structure in which a 3-fold axis was overlooked. They computed the refinement and determined
that Te4 Pt3 possesses a trigonal structure with R3m space group and lattice parameters a = 3.988
˚ and c = 35.39 A.
˚ Preliminary results of XRD and neutron diffraction experiments are summaA
rized in Appendix A. Therefore, it is necessary to check if any of these possible structures match
the diffraction pattern obtained from Ti4 Pt3 .
37
CHAPTER 4
Ti-Pt PHASE DIAGRAM
In light of the poor understanding of the Ti-Pt phase diagram in the composition range 30-50
at.% Pt, several alloys in this range were prepared and analyzed in the as-cast and heat treated
conditions. The purpose was to verify if the microstructures are in agreement with the Biggs’
phase diagram and to confirm the existence of the peritectoid reaction proposed by Biggs and
coworkers [1].
4.1
Experimental Methodology
Five alloys, namely Ti-31Pt, 35Pt, 39Pt, 42Pt and 44Pt (at.% Pt) were prepared at the NASA
Glenn Research Center by non-consumable arc melting high purity Ti (99.995 wt.%) and Pt
(99.995 wt.%) together in the appropriate proportions. Since platinum (21.45 g/cm3 ) is almost
five times denser than titanium (4.5 g/cm3 ) and has a slightly higher melting temperature, the buttons were flipped and remelted several times in order to insure complete melting and mixing. In
spite of these efforts to insure mixing, problems were occasionally encountered in the form of
unmelted Pt and this was taken into account in the different analyses.
4.1.1
Differential Thermal Analysis (DTA) and Heat Treatment
DTA was performed on four of the six alloys in a Netzsch Differential Scanning Calorimeter
(DSC 404 C) and in a Netzsch Simultaneous Thermal Analysis (STA 409 C), both using the heat
flux Differential Scanning Calorimetry mode (HF-DSC). Both instruments consist of a single furnace and two crucibles with thermocouples. One crucible is for the sample being tested and the
other is for a reference material, in this case, an empty Al2 O3 crucible. The mass of the samples
ranged from approximately 30 to 200 mg. The samples were heated and cooled twice from room
38
temperature to 1300 ◦C at a rate of 20 ◦C/min in an alumina crucible. The furnace was evacuated
and purged with Argon gas three times before the experiments to minimize contamination effects.
Finally, a continuous flow of argon of 64 mL/min was maintained during the experiments.
The results were used to design the heat treatments and verify the presence of the phase
transformations described by Murray [3] and Biggs et al. [1] in the composition range of interest.
The results of the DTA experiments for the second heating/cooling cycle are shown in Figure 4.1.
According to the experiments, a peak between 1220 and 1235 ◦C is observed during heating in
all the curves with variations in intensity. This peak falls in the same temperature range observed
by Biggs et al. [1] that resulted in the proposition of a peritectoid transformation by these authors.
A lower temperature peak between 930 and 1050 ◦C is observed in all the alloys on heating, but
in the low platinum alloys, this peak has an intensity much lower compared to the 39Pt and 44Pt
alloys. This peak falls in the temperature range of the martensitic transformation of the TiPt phase
(B19B2).
The DTA scan on the Ti-44Pt alloy revealed the presence of two additional peaks on cooling,
compared with the other alloys, at 1014 ◦C and 963 ◦C but no additional peaks were observed
on heating within the same temperature range. Biggs et al. [48] observed two additional peaks
on cooling and one additional on heating in alloys 43 and 50 at.% Pt. They explained that these
additional peaks corresponded to intermediate phases between the β-TiPt (B2) and α-TiPt (B19)
phases. Later, it will be shown that there could be an alternative explanation for these peaks.
Based on these results, the alloys were heat treated at three different temperatures, namely,
1260 ◦C, 1050 ◦C, and 800 ◦C to bracket the peaks observed during the DTA experiments (see
Figure 4.2a(a)) and, consequently, to verify the correctness of the phase diagram proposed by
Biggs et al. [1]. Prior to heat treatment, the samples were cleaned, wrapped in Ta foil and encapsulated in quartz ampoules that were evacuated to less that 13 Pa (100 mtorr) and backfilled
with argon in order to minimize interstitial contamination; in this manner, any impurity-stabilized
phases would be minimized. Three sets of alloys were heated to 1260 ◦C, held for 2 h and then
39
2.8
DSC (mW/mg)
2.6
exo
2.4
1225.7 ◦ C
0.04 J/g
2.2
2.0
1.8
cooling
heating
1.6
1.4
1.2
900
950
1000
1050
1100
1150
1200
Ti-31Pt
1250
1.2
DSC (mW/mg)
1.1
1228.7 ◦ C
3.80 J/g
exo
1.0
0.9
0.8
0.7
0.6
cooling
heating
1009.5
-0.05 J/g
0.5
900
DSC (mW/mg)
0.6
950
Ti-35Pt
1050
1100
cooling
0.4
0.2
1041.6 ◦ C
0.61 J/g
heating
950
1000
exo
963.2 ◦ C
1050
1100
1150
1200
1250
1152.7 ◦ C
-9.79 J/g
cooling
0.6
939.2 ◦ C
1222.0 ◦ C
7.09 J/g
0.4
900
1250
Ti-39Pt
1014.5 ◦ C
-0.35 J/g
0.8
0.2
1200
1234.1 ◦ C
8.84 J/g
1146.3 ◦ C
-7.63 J/g
0.3
1.0
1150
934.2 ◦ C
-1.11 J/g
0.0
900
DSC (mW/mg)
1000
exo
0.5
0.1
1146.6 ◦ C
-3.86 J/g
◦C
1044.6 ◦ C
4.34J/g
heating
950
1000
1050
Ti-44Pt
1100
Temperature (◦ C)
1150
1200
1250
Figure 4.1: Differential thermal analysis results obtained for Ti-31Pt, 35Pt, 39Pt, and 44Pt alloys.
40
one set was water quenched. The remaining sets were cooled to 1050 ◦C, held for another 2 h then
the second set was water quenched. Finally, the remaining set was cooled to 800 ◦C, held for 2 h
and water quenched, as is shown in the schematic in Figure 4.2a. A nomenclature for each of these
heat treatments that will be used in this chapter is summarized in Table 4.1.
44
WQ
1050
1050 ◦ C
2h
min)
1260
1260 ◦ C
2h
800
HT 1
HT 2
800 ◦ C
2h
WQ
WQ
6 (K/
35
Temperature ◦ C
31
Temperature ◦ C
39
HT 3
time (h)
at.% Pt
(a)
(b)
Figure 4.2: (a) Ti-Pt phase diagram showing the alloys heat treated at three different temperatures
to bracket the peaks observed during the DTA experiments. (b) Schematic of the heat treatments
performed on the alloys.
Table 4.1: Nomenclature and description of the heat treatments performed on the alloys.
Nomenclature
Description
HT1
as-cast sample heated to 1260 ◦C for 2 h followed by water quenching.
HT2
as-cast sample heated to 1260 ◦C for 2 h, furnace cooled to 1050 ◦C and maintained at this temperature for 2 h followed by water quenching.
HT3
as-cast sample heated to 1260 ◦C for 2 h, furnace cooled to 1050 ◦C and maintained at this temperature for 2 h, furnace cooled to 800 ◦C and maintained at
this temperature for 2 h followed by water quenching.
HT4
as-cast sample heated to 1260 ◦C for 2 h followed by air cooling.
41
4.1.2
Scanning Electron Microscopy (SEM) and Energy Dispersive Spectroscopy (EDS)
Samples for metallographic and microstructural analysis in the SEM were sectioned from the
melted buttons and the heat treated samples using a low speed saw. The microstructure from the
bottom to the top of the melted buttons was considered for these analyses as Figure 4.3 shown.
Later, samples were prepared by standard metallographic techniques including grinding (240-320400-600 grit SiC papers), polishing (6-3-1 μm diamond suspension) and vibratory polishing that
was carried out overnight using a 0.05 μm colloidal silica suspension on a Buehler Vibromet 2.
top
sample
sample
bottom
Lateral view
Top view
Figure 4.3: Schematic of the location where the samples for metallographic and microstructural
analysis were obtained from the melted buttons.
All the SEM images presented in this thesis are backscattered electron images (BEIs) obtained
in a JEOL JSM-7000F Field Emission Scanning Electron Microscope (FESEM) using 15 kV acceleration voltage and a medium probe current. Semi-quantitative compositional information on
the different phases was obtained using an EDAX Genesis Energy Dispersive Spectroscopy (EDS)
system. All the compositions measured were normalized using as a standard the composition of the
Ti-35Pt alloy measured at NASA Glenn Research Center summarized using Inductively Coupled
Plasma Atomic Emission Spectroscopy (ICP-AES) (Table 4.2).
Table 4.2: Chemical composition of the as-cast Ti-35Pt alloy measured at NASA Glenn Research
Center using ICP-AES.
Sample
Ti-35Pt
wt.%
at.%
Ti
Pt
O
N
C
Ti
Pt
31.1
68.8
0.0056
0.019
0.0033
64.8
35.2
42
4.1.3
Transmission Electron Microscopy (TEM)
Samples for TEM analysis were sectioned with a low speed saw, thinned using 240-320400-600 grit SiC papers. Later, discs of 3 mm in diameter, with average thickness of 150 μm,
were obtained using an abrasive cutting machine. Then, the discs were dimpled on a VCR D500i
Dimpler, and finally ion-milled on a Gatan Model 600 Duomill. Micrographs and selected area
diffraction patterns (SADPs) were obtained using both a Phillips CM12 operated at 120 kV accelerating voltage and a Phillips CM200 operated at 200 kV accelerating voltage. Compositional
information was obtained using an EDAX Genesis EDS system on the CM12 microscope, and a
PGT Prism EDS system on the CM200 microscope.
4.2
As-cast Microstructures
4.2.1
Hypoeutectic Alloy: Ti-31Pt
Backscattered electron images (BEIs) of the as-cast Ti-31Pt alloy (Figure 4.4a) indicate that
this alloy is hypoeutectic and consists of primary dendrites of Ti3 Pt and interdendritic eutectic.
The compositions of the phases taken on the FESEM are summarized in Table 4.3. Further studies
in the TEM (Figure 4.4b) revealed that the “eutectic” is complex and composed of Ti3 Pt (a partial
view of the dendrite is observed in the bottom left part of the image), α-TiPt (that transforms on
cooling from β-TiPt), a faulted phase that confirms the presence of Ti4 Pt3 in the binary system,
and a small globular particle (white arrows) whose stoichiometry is close to “Ti4 Pt” (see Chapter 3
for the characterization of this phase).
Table 4.3: Platinum content of the phases observed in the as-cast Ti-31Pt alloy determined by EDS
in the FESEM.
alloy
Ti-31Pt
Phase at.% Pt
Ti3 Pt
interdendritic
29.0
44.0
43
It is important to note that the composition of the interdendritic region corresponds to an
average composition of the phases encountered in that region and not to a particular phase. The
composition of the phases observed in Figure 4.4b was measured in the TEM and the EDS spectrum is presented in Figure 4.4c. The spectrum reveals that the particles of Ti4 Pt contain more
titanium than the other phases (Ti/Pt∼3.9) and more oxygen, suggesting that this is an oxygen
stabilized phase. The sequence of transformations upon cooling is as follow:
Ti4 Pt3 +
α-TiPt
Ti3 Pt
(a) 1,000X
Figure 4.4: Ti-31Pt alloy in the as-cast condition. (a) BEI. (b) TEM image from a region of the
eutectic observed in the BEI and a portion of Ti3 Pt dendrites (bottom left part of the image). The
eutectic is comprised of Ti4 Pt3 , α-TiPt and particles of Ti4 Pt (white arrows) in the interdendritic
regions. (c) The EDS spectra reveal that the particles of “Ti4 Pt” contain more titanium and slightly
more oxygen than the other phases present in the microstructure.
4.2.2
Eutectic Alloy: Ti-35Pt
The Ti-35Pt alloy appears to correspond to the eutectic composition as there was essentially
no primary phase in the BEIs. Significantly however, the higher magnification micrograph (Figure 4.5b) shows clearly the presence of three phases. The compositions of the observed phases,
44
1
31
11
1
22
0
α-TiPt
B=112particle
Ti4 Pt3
Ti3 Pt
(b) 28,000X
1.2
1.1
TiK
0.9
CK
0.3
NK
OK
PtM
1.0
0.7
0.2
Counts
0.8
TiK
0.5
0.1
0.3
0.0
0.0
0.5
1.0
0.1
4.4
4.5
4.6
0.6
0.4
Ti4 Pt3
Martensite
Ti3 Pt
Particle
PtL
0.2
PtL
TiK
PtM
PtL
PtL
PtM
PtL
0.0
0
2
4
6
8
Energy (keV)
(c)
45
10
12
14
namely Ti3 Pt, Ti4 Pt3 , and α-TiPt, are presented in Table 4.4. Assuming the high Z regions are indicative of the peritectoid transformation (β-TiPt+Ti3 Pt→Ti4 Pt3 ), it is clear from Figure 4.5b that
the peritectoid reaction does not go to completion, presumably because of sluggish transformation
kinetics, leaving untransformed β-TiPt, which in turn transforms to α-TiPt upon cooling through
the lower temperature DTA peak at approximately 930 ◦C.
Table 4.4: Pt content of the phases observed in the as-cast Ti-35Pt alloy as determined by EDS in
the FESEM.
Alloy
Ti-35Pt
Phase at.% Pt
Ti3 Pt
Ti4 Pt3
α-TiPt
29.0
41.7
44.0
(a) 1,000X
Figure 4.5: Ti-35Pt in the as-cast condition. It appears to correspond to the eutectic composition
because no primary phase is observed (a). The TEM image (b) reveals that the peritectoid transformation does not go to completion leaving β-TiPt that transforms into α-TiPt via the martensitic
transformation at around 930 ◦C.
46
Ti4 Pt3 +
α-TiPt
Ti3 Pt
(b) 5,000X
Ti4 Pt3
Ti3 Pt
α-TiPt
(c)
47
4.2.3
As-Cast Ti-39Pt Alloy
This alloy (Figure 4.6a) is a hypereutectic alloy consisting of dendrites of Ti4 Pt3 , confirmed
with TEM (Figures 4.6c and 4.6d), that apparently transformed from the primary β-TiPt dendrites
expected from the phase diagram, and eutectic in accordance with the phase diagram. A halo of
Ti3 Pt forms around the primary dendrites of β-TiPt separating the dendrites from the interdendritic
eutectic. Also, a layer of α-TiPt is observed near, but not at, the periphery of the primary dendrites
as well as in the high Z regions in the interdendritic eutectic (Figure 4.6b); as will be discussed
further below, this appears to be inconsistent with the peritectoid transformation proposed by Biggs
and coworkers. The Ti5 Pt3 was not possible to resolve in the FESEM, but TEM confirms its
presence (Figure 4.6c). The compositions of the phases are summarized in Table 4.5 along with
the compositions measured for the 42Pt and 44Pt alloys.
Table 4.5: Pt content of the phases observed in the as-cast Ti-39Pt, 42Pt and 44P alloys determined
by EDS in the FESEM.
Alloy
a
b
Phase, at.% Pt
Ti3 Pt
Ti5 Pt3
Ti4 Pt3
α-TiPt
Ti-39Pt
29.3
a
40.7
44.8
Ti-42Pt
29.8
b
40.8
45.3
Ti-44Pt
28.0
37.1
42.3
46.2
It was observed in the TEM.
The size of this phase was too small to measure by EDS in the FESEM.
4.2.4
The solidification structure of this alloy consisted of a large volume fraction of primary dendrites with the core comprised of a lamellar structure of Ti4 Pt3 and α-TiPt (which transforms from
β-TiPt upon cooling). The dendrites are surrounded by lower Z interdendritic regions (Figure 4.7a)
comprised of Ti3 Pt and Ti4 Pt3 phases. The high magnification BEI shown in Figure 4.7b revealed
48
Ti4 Pt3
(a) 1,000X
Ti4 Pt3
α-TiPt
Ti3 Pt
(b) 5,000X
Figure 4.6: Ti-39Pt alloy in the as-cast condition. (a) 1000X (b) 5000X. This is a hypereutectic
alloy with β-TiPt primary dendrites that transformed to Ti4 Pt3 upon cooling, and eutectic in the
interdendritic regions. The BEI in (b) reveals that the peritectoid reaction does not go to completion
in the interdendritic regions, and that a layer of α-TiPt phase is observed in the periphery of the
primary β-TiPt dendrites. The TEM images in (c) and (d) reveal the morphology of the phases
observed in the BEIs.
49
Ti4 Pt3
Ti4 Pt3
α-TiPt
Ti5 Pt3
(c) 45,000X
(d) 45,000X
that some of the interdendritic regions contain Ti5 Pt3 (see Chapter 3 for the characterization of
this phase) and small particles of a low Z phase within the Ti3 P and Ti4 Pt3 phases whereas some
other regions revealed a different microstructure consisting of Ti3 Pt and Ti5 Pt3 (Figure 4.7c). The
platinum content of the phases observed in the FESEM were determined by EDS with the exception of the low Z particles due to their small size. The results are summarized in Table 4.5.
TEM analysis revealed that the interdendritic regions contain Ti3 Pt, Ti5 Pt3 and small particles
of Ti4 Pt phase (Figures 4.8a and 4.8b). In addition, Figure 4.8a reveals the presence of dendrites
of Ti4 Pt3 with α-TiPt on the peripheries of the dendrites as can be observed in the BEIs. The
morphology of the Ti5 Pt3 phase shown Figure 4.7c can be correlated with the TEM image shown
in Figure 4.8b.
50
Ti4 Pt3
α-TiPt
Ti3 Pt+
low Z particles
(a) 1,000X
Ti4 Pt3
Ti4 Pt3 +Ti5 Pt3 +
low Z particles
α-TiPt
Ti3 Pt
(b) 5,000X
Figure 4.7: As-cast Ti-42Pt alloy. (a) Low magnification BEI reveals the presence of large dendrites of Ti4 Pt3 phase containing the lamellar structure of Ti4 Pt3 +α-TiPt in its core. The high
magnification in (b) reveals that the interdendritic regions are comprised of Ti3 Pt, Ti5 Pt3 , Ti4 Pt3
and small particles of a low Z phase. (c) Other locations of the interdendritic regions contain a
larger volume of the Ti5 Pt3 phase than the region in Figure 4.7b.
51
α-TiPt
Ti4 Pt3
Ti3 Pt
Ti5 Pt3
(c)
Ti4 Pt3
01
Ti3 Pt
01
1
10 0
0
Ti3 Pt
Ti4 Pt
Ti5 Pt3
α-TiPt
202 222
020
(a) 35,000X
(b) 75,000X
Figure 4.8: As-cast Ti-42Pt alloy. The TEM images revealed that the small particles observed in
the BEIs are the Ti4 Pt phase and that the elongated features in (b) are the Ti5 Pt3 phase. In (b) ,
from top to bottom, 001 SADP from Ti3 Pt, 1011 CBP from Ti5 Pt3 and 101 CBP from the low
Z Ti4 Pt phase.
52
4.2.5
The Ti-44Pt is a hypereutectic alloy with a considerably larger volume fraction of primary
dendrites (Figure 4.9a). The primary dendrites of β-TiPt have a platinum concentration gradient
along the dendrites that gives rise to different phases and morphologies. The Pt-rich β-TiPt in the
core transforms to α-TiPt at the martensitic transformation temperature whereas the region next
to the core transforms to the lamellar-type structure of Ti4 Pt3 +β-TiPt (then β-TiPt αTiPt on
cooling). Finally, the periphery of the dendrites transforms mostly to Ti4 Pt3 via the peritectoid
reaction and/or by the direct transformation from β-TiPt (discussed below). The interdendritic
regions are composed of a divorced eutectic comprising Ti3 Pt and the Ti5 Pt3 phases.
Ti4 Pt3
Ti3 Pt
Ti5 Pt3
α-TiPt
(a) 2,000X
Figure 4.9: The microstructure of as-cast Ti-44Pt alloy reveals the existence of a large volume
fraction of primary dendrites and divorced eutectic in the interdendritic regions. It is observed that
the cored primary dendrites of β-TiPt transformed into α-TiPt in the center, lamellar Ti4 Pt3 +βTiPt and Ti4 Pt3 on the edges upon cooling. In addition, Ti3 Pt and Ti5 Pt3 are observed in the
interdendritic regions where the impurity content is believed to be sufficient to stabilize the Ti5 Pt3
phase in this region. Some particles are observed in between the interdendritic regions and the
dendrites. The SADP pattern in (b) confirms that they are the Ti3 Pt phase.
53
Ti3 Pt
Ti4 Pt3
101 121
020
B= 101Ti3 Pt
(b) 35,000X
During solidification, solute and impurity atoms, such as oxygen, are rejected into the remaining liquid until the eutectic composition is reached. The impurity content is sufficient to stabilize
the Ti5 Pt3 phase in this region. This interdendritic region is different from the interdendritic regions observed in the microstructures of as-cast 31Pt, 35Pt and 39Pt alloys in which the eutectic
consists of Ti3 Pt, Ti4 Pt3 and β-TiPt (later, it transforms to α-TiPt during cooling). In addition,
some particles are observed in between the interdendritic region and the dendrites; these particles
were examined in the TEM (Figure 4.9b) and the SADP pattern confirmed that they are Ti3 Pt. The
compositions of the various phases in this alloy are summarized in Table 4.5.
4.3
Microstructures After Heat Treatment
4.3.1
Ti-31Pt Alloy
BEIs of the Ti-31Pt alloy at the three temperatures selected for heat treatment (Figures 4.10a
to 4.10c) reveal the presence of an additional intermediate Z phase, Ti5 Pt3 , in the interdendritic
54
regions. At 1260 ◦C (HT1), the bright phase corresponds to α-TiPt which transforms from β-TiPt
upon quenching. On the other hand, at 1050 ◦C (HT2) and 800 ◦C (HT3) the bright phase corresponds to Ti4 Pt3 . The compositions of the observed phases, determined by EDS in the FESEM, at
the three different temperatures are summarized in Table 4.6.
Table 4.6: Platinum content of the phases observed in the Ti-31Pt alloys heat treated at 1260 ◦C,
1050 ◦C, and 800 ◦C.
Temperature
Phase, at.% Pt
Ti3 Pt
Ti5 Pt3
Ti4 Pt3
α-TiPt
30.5
37.3
-
44.5
1050 C
30.5
37.5
42.2
-
800 ◦C
30.7
37.4
42.4
-
1260 ◦C
◦
α-TiPt
Ti3 Pt
Ti5 Pt3
(a) 1260 ◦C (HT1). 1,000X
Figure 4.10: Ti-31Pt alloy heat treated at (a) 1260 ◦C (HT1), (b) 1050 ◦C (HT2), and (c)
800 ◦C (HT3). Note the presence of the Ti5 Pt3 phase after these heat treatments compared with the
as-cast materials in Figure 4.4a.
55
Ti4 Pt3
Ti3 Pt
Ti5 Pt3
(b) 1050 ◦C (HT2). 1,000X
Ti3 Pt
Ti5 Pt3
Ti4 Pt3
(c) 800 ◦C (HT3). 1,000X
56
4.3.2
Ti-35Pt Alloy
BEIs of the Ti-35Pt alloy at the three temperatures (Figures 4.11a to 4.11c) also reveal the
presence of the intermediate Z Ti5 Pt3 phase. As in the Ti-31Pt alloy at 1260 ◦C (HT1), the
bright phase corresponds to α-TiPt which transformed from β-TiPt upon quenching, while at
1050 ◦C (HT2) and 800 ◦C (HT3) the bright phase corresponds to the Ti4 Pt3 phase. The compositions of the observed phases, determined by EDS in the FESEM, at these three different temperatures are summarized in Table 4.7.
1050 ◦C, and 800 ◦C.
Temperature
Phase, at.% Pt
Ti3 Pt
Ti5 Pt3
Ti4 Pt3
α-TiPt
1260 ◦C
30.5
37.3
-
44.5
1050 ◦C
30.4
37.6
41.8
-
30.8
37.3
42.3
-
◦
800 C
α-TiPt
Ti3 Pt
Ti5 Pt3
(a) 1260 ◦C, 2,000X
Figure 4.11: Ti-35Pt alloy heat treated at (a) 1260 ◦C (HT1), (b) 1050 ◦C (HT2) and 800 ◦C (HT3).
57
Ti4 Pt3
Ti3 Pt
Ti5 Pt3
(b) 1050 ◦C, 2,000X
Ti4 Pt3
Ti3 Pt
Ti5 Pt3
(c) 800 ◦C, 2,000X
4.3.3
Ti-39Pt Alloy
The microstructures of the Ti-39Pt alloy at the three temperatures (Figures 4.12 and 4.13) also
contain Ti5 Pt3 while the remaining phases are in agreement with the Biggs’ phase diagram. At
58
1260 ◦C (HT1), the dendrites are β-TiPt which later transformed to α-TiPt upon quenching (confirmed by TEM analysis, Figure 4.12b). The SADPs for Ti3 Pt, B=001, and for Ti5 Pt3 , B=2113,
are presented for confirmation. At 1050 ◦C (HT2) and 800 ◦C (HT3), the dendrites of β-TiPt transformed into Ti4 Pt3 . This was confirmed in the TEM (Figures 4.13b and 4.13d). The compositions
of the observed phases, determined by EDS in the FESEM, at the three different temperatures are
summarized in Table 4.8.
1050 ◦C, and 800 ◦C.
Temperature
Phase at.% Pt
Ti3 Pt
Ti5 Pt3
Ti4 Pt3
α-TiPt
1260 ◦C
31.0
37.5
-
43.9
1050 ◦C
30.5
37.2
42.2
29.9
36.7
42.8
◦
800 C
11 0 1 1
Ti5 Pt3
0110
011
Ti3 Pt
B= 2113Ti5 Pt3
α-TiPt
Ti3 Pt
α-TiPt
01
01
10
Ti5 Pt3
10
0
B= 001Ti3 Pt
(a) 1260 ◦C. 2,000X
(b) 1260 ◦C. 35,000X
Figure 4.12: (a) Ti-39Pt alloy heat treated at 1260 ◦C (HT1). The microstructure is composed of
dendrites of β-TiPt that later transform to α-TiPt upon quenching, Ti3 Pt and Ti5 Pt3 . SADPs from
Ti3 Pt, B=001, and Ti5 Pt, B=2113 are presented for confirmation in (b).
59
Ti4 Pt3
201
211
010
B= 102Ti3 Pt
Ti3 Pt
Ti3 Pt
Ti4 Pt3
Ti5 Pt3
1010
Ti5 Pt3
B= 0001Ti5 Pt3
(a) 1050 ◦C. 2,000X
(b) 1050 ◦C. 28,000X
0
01
20
12
11
Ti3 Pt
Ti3 Pt
B= 102Ti3 Pt
Ti5 Pt3
Ti4 Pt3
Ti4 Pt3
B= 2110Ti5 Pt3
0110
Ti5 Pt3
(c) 800 ◦C. 2,000X
0112
0002
(d) 800 ◦C. 60,000X
Figure 4.13: (a) Ti-39Pt alloy heat treated at 1050 ◦C (HT2). Three phase are present in the
microstructure Ti4 Pt3 , Ti3 Pt, and Ti5 Pt3 . SADPs for Ti3 Pt, B=102, and Ti5 Pt3 , B=0001, are
presented for confirmation in (b). In the case of the SADP from Ti3 Pt, the 001 reflection is present
due to double diffraction. (c) Ti-39Pt alloy heat treated at 800 ◦C (HT3). Three phase are present
in the microstructure Ti4 Pt3 , Ti3 Pt, and Ti5 Pt3 . SADPs for Ti3 Pt, B=102, and Ti5 Pt, B=2110,
are presented for confirmation in (d). In the case of the SADP from Ti3 Pt, the 001 reflection is
present due to double diffraction.
60
4.3.4
Ti-44Pt Alloy
The microstructures of the Ti-44Pt alloy heat treated at the three temperatures are presented in
Figures 4.14 and 4.15. At 1260 ◦C (HT1), only two phases were observed (Figure 4.14a), namely,
Ti5 Pt3 and α-TiPt. Their presence was confirmed in the TEM (Figure 4.14b). An SADP from
Ti5 Pt3 , B=2110, is presented for confirmation. This observation is inconsistent with the Biggs’
phase diagram (Figure 2.3c).
0002 0112
0110
Ti5 Pt3
B= 2110Ti5 Pt3
α-TiPt
α-TiPt
Ti5 Pt3
(a) 1,000X
(b) 60,000X
Figure 4.14: Ti-44Pt alloy heat treated at 1260 ◦C (HT1). The microstructure reveals the presence
of two phases β-TiPt (it transforms to α-TiPt upon quenching) and Ti5 Pt3 . The SADP from Ti5 Pt3 ,
B=2110, is presented for confirmation in (b).
At 1050 ◦C (HT2), the Ti-44Pt microstructure reveals the presence of four phases (Figure 4.15a).
The prior β-TiPt dendrites transform into a lamellar-type structure of β-TiPt and Ti4 Pt3 , in which
the former transforms into α-TiPt upon quenching. In addition, the interdendritic regions consist
of two low Z phases. Detailed analysis in the TEM revealed that these phases were Ti5 Pt3 (dark
contrast in Figure 4.15a) and Ti4 Pt3 (medium dark contrast in Figure 4.15a). The latter Ti4 Pt3 contains a slightly higher Ti content compared to the surrounding Ti4 Pt3 , as is shown clearly in the
inset of the TiKα peak in the range 4.45-4.60 keV in Figure 4.15d. This will be discussed further
61
in Section 4.2.
Ti4 Pt3 (I)
α-TiPt
Ti5 Pt3
Ti5 Pt3
Ti4 Pt3 (II)
Ti4 Pt3 (II)
1101 0111
1211
Ti4 Pt3 (I)
B= 1011Ti5 Pt3
(a) 1,000X
(b) 60,000X
PtM
1.0
Ti4 Pt3
0.7
0.8
Counts
TiK
Ti4 Pt3 (II)
0.5
Ti4 Pt3 (I)
0.4
4.45
4.50
4.55
4.60
4.65
0.6
0.4
α-TiPt
Ti5 Pt3
0.6
Ti5 Pt3
Ti4 Pt3 (II)
Ti4 Pt3 (I)
PtL
0.2
PtM PtM
TiK
PtL
PtL
PtL
PtL
0.0
0
2
(c) 45,000X
4
6
8
Energy (keV)
10
12
14
(d)
Figure 4.15: Ti-44Pt alloy heat treated at 1050 ◦C (HT2). (a) BEI reveals the presecence of four
phases: α-TiPt, Ti4 Pt3 , Ti5 Pt3 and a medium gray phase. In (b) TEM image of this medium
gray phase indicates that it is Ti4 Pt3 and the EDS spectra in (d) obtained from the phases shown in
Figure 4.15b revealed that this Ti4 Pt3 (II) adjacent to the Ti5 Pt3 contains slightly more Ti compared
to the matrix Ti4 Pt3 (I).
At 800 ◦C (HT3), Figure 4.16, the microstructure reveals that the prior β-TiPt dendrites transform to a lamellar-type structure of Ti4 Pt3 (gray region) and a higher Z region that, at closer exam-
62
ination, reveals a fine lamellar structure. Detailed analysis in the TEM (Figures 4.16c and 4.16d)
showed that this fine lamellar structure is comprised of Ti4 Pt3 and α-TiPt where, unlike the other αTiPt regions observed in the previous as-cast and heat treated alloys, does not form martensitically.
The low Z interdendritic regions in the BEIs are composed of Ti5 Pt3 (dark contrast in Figure 4.16a)
and Ti4 Pt3 (II) (medium dark contrast in Figure 4.16a). As in the previous sample, the latter Ti4 Pt3
contains a higher Ti content compared to the surrounding Ti4 Pt3 (I) (light gray phase). The platinum
content of the phases observed at these three different temperatures are summarized in Table 4.9.
It was not possible to determine the platinum content of the Ti4 Pt3 phase inside the fine lamellar
structure by EDS in the FESEM because of the small size of the phase.
1050 ◦C, and 800 ◦C.
Temperature
◦
4.4
Phase at.% Pt
C
Ti3 Pt
Ti5 Pt3
Ti4 Pt3 (I)
Ti4 Pt3 (II)
α-TiPt
1260
-
38.0
-
-
44.0
1050
-
38.0
42.3
41.5
48.1
800
-
37.5
42.6
41.3
50.2
Discussion
4.4.1
Transformations Observed in the as-cast Alloys
Clearly, differences exist between the microstructures observed and the predictions of the
Biggs’ phase diagram and these will be discussed further below. Before continuing with the discussion, it is important to clarify that the α-TiPt phase observed in the as-cast alloys had transformed
from β-TiPt upon quenching. Having said that, the microstructure of the as-cast Ti-31Pt alloy reveals that a lamellar structure of Ti4 Pt3 and α-TiPt exists in the interdendritic regions (resolved in
the TEM image shown in Figure 4.4b). This means that β-TiPt can follow a transformation path
different from the peritectoid transformation suggested by Biggs et al. [1].
63
α-TiPt+
Ti4 Pt3
Ti4 Pt3 (I)
Ti5 Pt3
α-TiPt+
Ti4 Pt3
Ti4 Pt3
Ti4 Pt3 (II)
(a) 2,000X
(b) 5,000
B= 001α-TiPt
20
α-TiPt
0
21
0
01
0
α-TiPt
Ti4 Pt3
Ti4 Pt3
(c) 35,000
(d) 75,000
◦
Figure 4.16: Ti-44Pt alloy heat treated at 800 C (HT3). (a) BEIs revealed the presence of Ti5 Pt3 ,
α-TiPt, Ti4 Pt3 (I) and Ti4 Pt3 (II); the latter two observed also in the Ti-44Pt alloy subjected to HT2.
The BEI in (b) and the TEM images in (c-d) revealed the presence of a fine lamellar structure of
Ti4 Pt3 and α-TiPt that had transformed from the prior β-TiPt lamellae observed at 1050 ◦C.
The as-cast Ti-35Pt alloy corresponds to the eutectic composition. The microstructure contains three phases Ti3 Pt, Ti4 Pt3 and α-TiPt and, at first sight, it is possible to conclude that the
peritectoid reaction is taking place but is limited to the Ti3 Pt/β-TiPt regions presumably due to
the sluggish nature of the transformation. It is possible to observe the peritectoid transformation
in several places of the microstructure (red circles in Figure 4.17); however, it is also possible to
64
observe a lamellar structure of Ti4 Pt3 and α-TiPt in other regions of the microstructure (red square
in Figure 4.17).
Ti4 Pt3 +
α-TiPt
Ti3 Pt
lamellar
structure of
Ti4 Pt3 + α-TiPt
peritectoid
transformation
Figure 4.17: BEI of the Ti-35Pt as-cast alloy. Some regions of the microstructure appear to be
consistent with the peritectoid transformation (red circles) whereas other larger high Z areas form
a lamellar structure suggesting the precipitation of Ti4 Pt3 directly from the β-TiPt phase with no
involvement of the Ti3 Pt phase (red square).
In the as-cast Ti-39Pt alloy, a layer of α-TiPt exists near the periphery of the β-TiPt dendrites
which transform to Ti4 Pt3 upon cooling, Figure 4.6b. This layer is separated from the interdendritic
Ti3 Pt phase by a layer of Ti4 Pt3 . This observation suggests that the solvus line on the Ti-rich side
of the β-TiPt phase is incorrect. An EDS line scan of the dendrites (46 points along the line
shown in Figure 4.18a) revealed that the Pt concentration remains approximately constant along
its cross-section (Figure 4.18b), with the exception of the peripheries where the Pt-rich α-TiPt
phase exists. Considering 44 points as the size of the data, not considering the periphery of the
dendrites, the mean platinum composition (M ) within the dendrites is approximately 41.5 at.% Pt,
which is close to the composition of the Ti4 Pt3 phase as measured by Biggs et al. [1]. The standard
deviation (SD) from the mean is 0.92.
This result suggests that the β-TiPt solidus line should be displaced to lower Pt and, also, it
should be relatively vertical over the solidification temperature range of this particular alloy such
that little coring accompanies solidification. Likewise, the β-TiPt solvus must exhibit decreasing
65
60
56
52
48
at.%
44
40
36
line scan
32
Pt
M=41.5 at.% Pt
SD=±0.92
28
24
20
0
4
(a)
8
12
16
20
24
28
32
Number of points
36
40
44
48
(b)
Figure 4.18: (a) BSE image indicating the location of the EDS line scan. (b) Results of the line
scan. Note that the platinum concentrations remain approximately constant along the cross section
of the dendrites. The mean Pt composition is 41.5 at.% Pt and the standard deviation (SD) from
the mean is 0.92.
Ti solubility (increasing Pt) until the original solvus line is reached at ∼ 1100 ◦C, which takes
into consideration the composition of the α-TiPt measured at the three different temperatures of
the heat treatments in the Ti-39Pt and Ti-44Pt alloys (Tables 4.8-4.9). This proposed modification
(Figure 4.19) is corroborated by the differences in the volume fraction of dendrites measured experimentally (using ImageJ software [49, 50]) and the volume fraction calculated using the Scheil
Equation 4.1:
fs = 1 −
Cs
kCo
1
1−k
(4.1)
where fs is the solid fraction, in this case the volume fraction of the primary dendrites, CS the
composition of the solid, Co the composition of the alloy, k the partition coefficient given by
k = CS /CL , and CL the composition of the liquid. Considering CS = 46, Co = 39 and k =
46/35 = 1.31 (at the eutectic temperature) from the Biggs’ diagram, the calculated solid fraction
is 29%, while the experimental value obtained with ImageJ [49] is approximately 50% (three BEIs
at 100X, 200X, 500X and 1000X were considered to calculate the solid fraction in ImageJ. The
66
mean value of all these measurements is given here with SD equal to 2.3%). On the other hand,
considering the modified diagram (Figure 4.19), the calculated fraction of primary dendrites using
CS = 41.5, Co = 39 and k = 41.5/35 = 1.12, is approximately 43%, which is closer to the 50%
observed empirically.
41.5 Pt
Figure 4.19: Modification to the Ti-Pt phase diagram appears as the dashed lines for the β-solidus
and the β-solvus.
The modified phase diagram in Figure 4.19 can also be used to explain the existence of the αTiPt layer observed in the as-cast Ti-39Pt alloy and an attempt to explain this is provided schematically in Figures 4.20 and 4.21. The following analysis considers that the transformations begin at
x = 0 (located at the initial β-TiPt/Ti3 Pt interface) and extend towards the core of the dendrite;
in addition, a planar Ti3 Pt/β-TiPt is considered (Figure 4.21). During solidification, β-TiPt dendrites solidify first with an essentially constant composition of 41.5 at.% Pt (Figure 4.20a). At
1424 ◦C the eutectic reaction takes place and Ti3 Pt nucleates and grows at the β-TiPt interface
and then the eutectic structure develops from the remaining liquid (Figures 4.20b and 4.21a). Between 1424 ◦C and 1205 ◦C, the composition of the β-TiPt should increase gradually from 41.5
67
to 45 at.% Pt which is achieved by the transfer of atoms across the β-TiPt/Ti3 Pt interface; this,
in turn, produce a local increase in the volume fraction of the Ti3 Pt phase. This is achieved by
the movement of the β-TiPt/Ti3 Pt interface to the left in Figures 4.20c and 4.21b. At 1205 ◦C,
the peritectoid reaction takes place at the β-TiPt/Ti3 Pt interfaces and, because the transformation
is sluggish, the reaction does not go to completion leaving behind a layer of Ti4 Pt3 in the edges
of the dendrites (Figures 4.20d and 4.21c). Below 1205 ◦C, the composition of the dendrites far
away from the interface favors the transformation of β-TiPt directly to Ti4 Pt3 (Figure 4.21d); near
the interface however, the Pt-rich β-TiPt does not transform until the martensitic temperature is
reached, accounting for the layer of α-TiPt near the dendrite edges (Figures 4.20e and 4.21e).
44
48
β-TiPt
Liquid
44
at.% Pt
40
39
36
36
32
28
28
40
(a) T > 1424◦C
48
45
44
β-TiPt
30
(b) T = 1424◦C
Ti3 Pt
distance from the core
of the dendrite
(c) 1205 ◦C < T < 1424◦C
α-TiPt
48
45
44
41.5
Ti4 Pt3
at.% Pt
at.% Pt
28
of the dendrite
36
32
40 Ti4 Pt3
30
28
of the dendrite
(d) T = 1205 ◦C
Ti4 Pt3
36
32
30
28
movement of
the interface
32
41.5
40
transfer of
atoms
36
30
of the dendrite
Ti3 Pt
initial interface
position
41.5
40
32
β-TiPt
45
44
41.5
41.5
at.% Pt
48
Ti3 Pt
β-TiPt
at.% Pt
48
Ti3 Pt
of the dendrite
(e) T < 1050 ◦C
Figure 4.20: Schematic drawings indicating the sequence of events taking place during solidification that lead to the formation of the α-TiPt layer that forms near the periphery of the dendrites in
the 39Pt alloy. See text for details of the proposed mechanism.
The direct transformation of the β-TiPt dendrites to Ti4 Pt3 can be explained using the schematic
68
β-TiPt
β-TiPt
Ti3 Pt
Ti3 Pt
0
x
(a) T = 1424◦C
x
0
(b) 1205 ◦C < T < 1424◦C
Ti4 Pt3
0
(c) T = 1205 ◦C
Ti4 Pt3
Ti4 Pt3
β-TiPt
Ti3 Pt
β-TiPt enriched
at the interface
initial position of β-TiPt enriched
of the interface
at the interface
x
β-TiPt
Ti3 Pt
Ti3 Pt
β-TiPt enriched
at the interface
α-TiPt
growth
Ti4 Pt3
x
◦
0
x
◦
(d) 1050 C < T < 1205 C
0
(e) T < 1050 ◦C
Figure 4.21: Schematic sequence of transformations in the Ti-39Pt alloy during solidification. The
x-axis has its origin at the β-TiPt/Ti3 Pt interface and extends towards the core of the dendrites.
Gibbs free energy curves presented in Figures 4.22 and 4.23. At T = Tp − ΔT the β-TiPt (Ptenriched located at the peripheries of the dendrites as shown in Figure 4.20c) and Ti3 Pt phase have
been undercooled (ΔT ) below the peritectoid temperature (Tp ); therefore, there is a driving force
(ΔGp ) available to nucleate and grow the Ti4 Pt3 phase. The peritectoid reaction takes place between Ti3 Pt of ∼ 30 at.% Pt and β-TiPt of ∼ 46 at.% Pt producing, as a result, Ti4 Pt3 containing
∼ 42 at.% Pt at the interfaces (dashed line 1) (Figure 4.22).
The peritectoid reaction does not go to completion and the dendrites far away from the interface still contain around 41.5 at.% Pt (see Figure 4.20d). At a temperature below the peritectoid
(Tp ), if a tangent is drawn to the Xd (41.5 Pt) composition on the Gibbs free energy curve of the βTiPt phase, this composition falls below the X(To ). Under this condition, thermodynamics allows
the partitionless transformation of the β-TiPt phase to Ti4 Pt3 both of Xd Pt. The driving force for
this transformation is ΔGd (Figure 4.23).
The Ti-42Pt alloy contains a considerably larger volume fraction of primary dendrites, com-
69
G
T = Tp − ΔT
Ti3 Pt
Ti4 Pt3
β-TiPt
ΔGp
1
40 42 44 46 at.% Pt
30
Figure 4.22: Schematic free energy curves for the transformations in the Ti-39Pt alloy during solidification. The peritectoid transformation takes place at the peritectoid temperature (Tp ) involving
Ti3 Pt of ∼ 30 at.% Pt and β-TiPt of ∼ 46 at.% Pt, reacting to form Ti4 Pt3 of ∼ 42 at.% Pt.
G
Ti3 Pt
T < TP
Ti4 Pt3
β-TiPt
X(To )
ΔGd
Xd X1
X2
40 42 44 46 at.% Pt
30
Figure 4.23: Schematic free energy curves for the transformations in the Ti-39Pt alloy during
solidification. At a temperature below the peritectoid temperature Tp , a β-TiPt of 41.5 at.% Pt can
transforms in a partitionless manner to Ti4 Pt3 of the same composition.
pared with the Ti-39Pt alloy, and a greater amount of microsegregation as reflected by the lamellar
structure of Ti4 Pt3 and α-TiPt in the core of the dendrites. The formation of this lamellar structure
will be explained in Section 4.4.2. It is interesting to note that two possible microstructures can be
obtained in the interdendritic regions of this alloy. The first contains Ti3 Pt, Ti4 Pt3 , Ti5 Pt3 and small
particles of the “Ti4 Pt” phase. The second contains mostly Ti3 Pt and Ti5 Pt3 . These differences
could be due to the amount of interstitial atoms rejected to the liquid during solidification due to
the low solubility of impurity atoms in the Ti4 Pt3 , Ti3 Pt and β-TiPt phases (β-TiPt → α-TiPt upon
cooling).
70
The as-cast microstructure of Ti-44Pt alloy revealed a large volume fraction of primary dendrites as the Ti-42Pt alloy but, in contrast, it revealed more microsegregation (Figure 4.24a).
Specifically, the primary dendrites of β-TiPt have a platinum content that decreases towards the
interdendritic regions, as is shown in Figure 4.24b. This concentration gradient along the dendrites gives rise to different phases and morphologies. The Pt-rich β-TiPt in the core transforms
to α-TiPt at the martensitic transformation temperature. Subsequently, the region next to the core
region (lower platinum content) transforms to the lamellar-type structure of Ti4 Pt3 +β-TiPt which
transforms to αTiPt on cooling. Finally, the periphery of the dendrites transform to Ti4 Pt3 via the
peritectoid reaction (Figure 4.22).
The interdendritic regions of the Ti-44Pt alloy are composed of a divorced eutectic comprising Ti3 Pt and what appears to be an intermediate phase of approximate composition Ti5 Pt3 . During
solidification, solute and impurity atoms, such as oxygen, are apparently rejected into the remaining liquid until its composition reaches the eutectic composition. If sufficient, the impurity content
appears to stabilize the Ti5 Pt3 phase in this region. This eutectic is different from the eutectic observed in the previous alloys (31Pt, 35Pt and 39Pt) that is comprised of Ti3 Pt, Ti4 Pt3 and α-TiPt.
In addition, some particles are observed in between the interdendritic regions and the dendrites
(white circle in Figure 4.24a). The particles were observed in the TEM (Figure 4.9b) and the
SADP pattern confirms that they are Ti3 Pt. The platinum content of the phases are summarized in
Table 4.5.
In this as-cast Ti-44Pt alloy, the peritectoid reaction does not take place because there is not
sufficient Ti3 Pt phase to participate in the transformation. The following analysis considers that
the transformations begin at x = 0 (located at the initial interdendritic/Ti3 Pt interface) and extend towards the core of the dendrite; in addition, a planar interdendritic/β-TiPt is considered (Figure 4.25). In region I, the Ti4 Pt3 phase transforms at the peripheries of the prior β-TiPt dendrites
via the partitionless transformation described in Figure 4.23 and grows towards the core of the
dendrites (Figure 4.25a) given that the composition in this region is ∼ 41.5 at.% Pt (Figure 4.24b).
According to the proposed modification of the Ti-Pt phase diagram shown in Figure 4.19, the
71
54
at.% Pt
Ti3 Pt
50
Ti4 Pt3 +
α-TiPt
Ti4 Pt3
Ti4 Pt3
Ti4 Pt3 +
α-TiPt
α-TiPt
Ti4 Pt3
at.% Pt
46
42
1 2 3 4 5 6 7 8 9 101112 131415 161718 19
+ + + + + + + + + + + + + + + ++ + +
38
α-TiPt
34
Ti5 Pt
30
0
(a)
2
4
6
8
10
12
Number of points
14
16
18
20
(b)
Figure 4.24: as-cast Ti-44Pt alloy. (a) BSE image indicating the location of the EDS line scan.
(b) Results of the line scan. Note that the platinum concentration decreases from the core of the
dendrite towards the periphery resulting in the different phases and morphologies observed in the
microstructure.
volume fraction of the Ti3 Pt phase increases in the temperature range between the eutectic and
peritectoid temperatures; therefore, the Ti3 Pt phase precipitates near the periphery of the dendrites
(Figure 4.24a).
The alloy continues cooling and given that the Pt content increases towards the core of the
dendrites, then, it is not longer possible to transform the β-TiPt phase in a partitionlessly manner
and, the interface Ti4 Pt3 /β-TiPt disrupts (region II). The composition in this region favors the decomposition of the β-TiPt phase into the lamellar Ti4 Pt3 +β-TiPt structure as shown in Figure 4.25b
(detailed explanation in Section 4.4.2). The core of the dendrites (region III) and the β-TiPt lamellae are enriched in platinum (∼ 46 at.% Pt) which at ∼ 1050 ◦C transform to α-TiPt via the
martensitic transformation (Figure 4.25c).
4.4.2
Transformation β-TiPt β-TiPt+Ti4 Pt3
The large peak observed around 1230 ◦C in the DTA experiments (Figure 4.1) suggests that
a significant enthalpy is involved in the transformation. However, the peritectoid reaction taking
72
II
I
interdendritic
region
α-TiPt
Ti4 Pt3
III
interdendritic
region
β-TiPt
interdendritic
region
growth
Ti4 Pt3
x
I
II
I
β-TiPt
Ti4 Pt3
Ti4 Pt3
0
(a) T1 < 1205◦C
x
x
0
x
III
0
(c) T < 1050◦C
(b) T2 < T1
II
I
0
(d)
Figure 4.25: Schematic sequence of transformations in the Ti-44Pt alloy during solidification. The
x-axis has its origin at the β-TiPt/Ti3 Pt interface and extends towards the core of the dendrites.
Regions I, II and III correlate with the as-cast microstructure presented in (d).
place along the Ti3 Pt/β-TiPt interfaces is expected to be sluggish in nature and, probably, can
not account for the amount of energy absorbed on heating or evolved on cooling. Therefore, a
different type of transformation must be taking place at this temperature. As shown below, the
transformation appears to correspond to the decomposition of the β-TiPt phase to Ti4 Pt3 +β-TiPt
(which transforms to α-TiPt upon cooling).
With the purpose of understanding the transformation taking place at around 1230 ◦C (Figure 4.26a), as-cast Ti-39Pt samples were heat treated to three different temperatures, namely
1220 ◦C, 1240 ◦C and 1260 ◦C, each one followed by an immediate water quench. The microstructure of the samples presented in Figures 4.26b to 4.26d reveals the presence of the mentioned lamellar-type structure of β-TiPt and Ti4 Pt3 within the dendrites. The lamellae are irregular
in nature. In addition, it seems that the amount of Ti3 Pt phase inside the dendrites increases gradually with increasing temperature. It is interesting to note that the microstructure of the sample
heat treated at 1260 ◦C was not 100% martensite as was the case for the sample heat treated at this
73
temperature for 2 h (HT1, Figure 4.12a). This could indicate that the samples cooled to below the
high temperature peak prior to quenching.
0.7
exo
α-TiPt+
g
coolin
Ti4 Pt3
DSC (mW/mg)
0.6
0.5
1234.1 ◦ C
8.84 J/g
Ti3 Pt
0.4
Ti5 Pt3
0.2
1100
1150
1200
Temperature (◦ C)
1260 ◦ C
g
heatin
1240 ◦ C
1220 ◦ C
0.3
1250
1300
(b) 1220 ◦C 0 h + WQ 2,000X
(a)
α-TiPt+
Ti4 Pt3
Ti3 Pt
Ti5 Pt3
(c) 1240 ◦C 0 h + WQ
(d) 1260 ◦C 0 h + WQ
Figure 4.26: Samples of the Ti-39Pt alloy were heated to 1220 ◦C, 1240 ◦C and 1260 ◦C followed
by an immediate WQ. (a) Heat treatments carried out to capture the transformation observed at
around 1230 ◦C in the DTA scans. The microstructures in (b-c) reveal the presence of a lamellar
structure of Ti4 Pt3 +β-TiPt (which transforms to α-TiPt upon quenching).
The formation of the lamellar β-TiPt+Ti4 Pt3 structure can be explained using the schematic
Gibbs free energy curves presented in Figures 4.27 and 4.28. At the peritectoid temperature Tp ,
the peritectoid reaction takes place between Ti3 Pt of ∼ 30 at.% Pt and β-TiPt of ∼ 46 at.% Pt, to
produce Ti4 Pt3 of ∼ 42 at.% Pt at the interfaces (dashed line 1 in Figure 4.27). The peritectoid
74
reaction does not go to completion and the remaining β-TiPt dendrites of ∼ 44 at.% Pt (Table 4.8).
The system will reduce the free energy (ΔGl ) by forming β-TiPt in Ti4 Pt3 of ∼ 42 at.% Pt and a
β-TiPt of ∼ 46 at.% Pt (Figure 4.28). On quenching, the β-TiPt transforms to martensite.
G
T = Tp − ΔT
Ti3 Pt
Ti4 Pt3
β-TiPt
ΔGp
1
40 42 44 46 at.% Pt
30
Figure 4.27: Schematic Gibbs free energy curves used to explain the sequence of transformations
in the Ti-39Pt alloy during heat treatment. First, the peritectoid transformation takes place at
T = Tp − ΔT (ΔT is the undercooling).
G
T < TP
Ti3 Pt
Ti4 Pt3
β-TiPt
X(To )
ΔGl
Xe
Xo
40
42 44 46 at.% Pt
30
Figure 4.28: Schematic Gibbs free energy curves used to explain the sequence of transformations
in the Ti-39Pt alloy during heat treatment. Second the transformation of β-TiPt to the lamellar
Ti4 Pt3 +β-TiPt structure.
To determine if the quenching step was carried out incorrectly, the heat treatment HT1 was
repeated (Table 4.1). It is observed in Figure 4.29a that some dendrites did transform to the lamellar β-TiPt+Ti4 Pt3 structure, while other dendrites transformed completely to α-TiPt (martensite).
Moreover, the peritectoid reaction takes place at all of the Ti3 Pt/β-TiPt interface as observed in
75
the high magnification image presented on Figure 4.29b. Therefore, these observations confirmed
that the quenching step was most likely carried out at at temperature below 1150 ◦C. Nevertheless,
they permit us to conclude that the transformation associated with the high temperature peak in the
DTA corresponds to β-TiPt β-TiPt+Ti4 Pt3 and results in the the lamellar structure observed in
Figures 4.26b to 4.26d, 4.29a and 4.29b.
Ti3 Pt
Ti4 Pt3 phase, product
of the peritectoid transformation
α-TiPt
Ti5 Pt3
(a) 1260 ◦C 2h + WQ. 2,000X
(b) 1260 ◦C 2h + WQ. 7,000X
Figure 4.29: Ti-39Pt alloy subjected to heat treatment HT1 (Table 4.1) (a) the microstructure reveals that some regions of the dendrites have transformed to 100% β-TiPt, while others have transformed to the lamellar Ti4 Pt3 +β-TiPt structure (then β-TiPt transforms to α-TiPt upon quenching).
(b) The peritectoid transformation can be observed along the interfaces Ti3 Pt/β-TiPt.
The heat treatment HT1 was repeated to replicate the microstructure encountered in Figure 4.12a.
The purpose was to verify if the microstructure shown in Figure 4.29a resulted from quenching below 1150 ◦C (below the high intensity DTA peak). The microstructure obtained was in agreement
with Figure 4.12a, and 100 % α-TiPt was obtained in the dendrites (Figure 4.30a). If the specimen is subjected to HT4 (Table 4.1), β-TiPt transformed to the lamellar Ti4 Pt3 +α-TiPt structure
(Figure 4.30b). This confirms that the DTA peak at high temperatures is associated with the transformation β-TiPt → β-TiPt+Ti4 Pt3 .
The same orientation relationship (OR) was determined for i) the Ti4 Pt3 phase that transforms
via the peritectoid reaction and α-TiPt that transforms from β-TiPt upon cooling (SADP in Fi-
76
Ti3 Pt
α-TiPt
Ti5 Pt3
(a) 1260 ◦C 2h + WQ. 2,000X
(b) 1260 ◦C 2h + air cooling. 2,000X
Figure 4.30: (c) The heat treatment HT1 was repeated resulting in 100% α-TiPt dendrites (βTiPt→ α-TiPt). (d) After air cooled a sample heat treated HT1 the lamella Ti4 Pt3 +β-TiPt structure
is obtained in the prior β-phase.
gure 4.31a), and ii) the Ti4 Pt3 obtained from the transformation of β-TiPt to the lamellar Ti4 Pt3 +βTiPt structure (SADP in Figure 4.31b). In both SADPs, the strong reflections corresponds to the
α-TiPt phase, while the weak reflections correspond to the Ti4 Pt3 phase. The OR can be expressed
by Equation 4.2:
0001Ti4 Pt3 110α−TiPt 111β−TiPt
(4.2)
{1010}Ti4 Pt3 (111)α−TiPt {110}β−TiPt
4.4.3
Transformation of Ti4 Pt3 from α-TiPt at low temperatures
The same procedure was carried out in the as-cast Ti-39Pt alloy but at a lower temperature, i.e,
this alloy was heat treated at 1050 ◦C followed by an immediate water quench. The microstructure
reveals that the layer of α-TiPt observed at the periphery of the dendrites disappeared and transformed to Ti4 Pt3 (Figures 4.32a and 4.32b). This suggests that the metastable α-TiPt transforms
to Ti4 Pt3 upon heating. To investigate this phenomenon, new DTA scans were run on the Ti-39Pt
77
111
111
1010
111
002
001
002
1010
111
(b)
(a)
Figure 4.31: SADPs reveal the OR between the Ti4 Pt3 and α-TiPt, which transforms from β-TiPt
upon cooling. (a) Transformation of β-TiPt to the lamellar Ti4 Pt3 +β-TiPt structure. (b) Peritectoid
reaction. Red fonts denote the reflections from the α-TiPt phase, while the blue fonts denote
the reflections from the Ti4 Pt3 phase. B= 0001 from Ti4 Pt3 (with pseudo-heaxoganl axes, see
Chapter 3 for the charaterization of this phase), and B= 110 from α-TiPt.
sample that was heated following HT1 treatment.
Samples containing 100% α-TiPt dendrites (Figures 4.12a and 4.30a), and those containing
the lamellar α-TiPt+Ti3 Pt structure in the dendrites (Figure 4.29a) were rescanned in the DTA.
Both samples were heated until 900 ◦C at a rate of 40 K/min at which point the heating rate was
reduced to 20 K/min. The former sample was cycled three times on heating and cooling between
900 ◦C and 1250 ◦C; while the latter sample was cycled twice on heating and cooling between
900 ◦C and 1070 ◦C.
The results of the DTA scans, presented in Figures 4.33a and 4.33b, reveal the existence of
an exothermic peak at around 690 ◦C on the first heating cycle for both samples. This implies that
the α-TiPt phase, obtained from β-TiPt phase during the water quenching, transforms to a more
stable lower temperature phase liberating heat in the process. From the BEIs above, this could be
associated with the transformation of α-TiPt phase into Ti4 Pt3 . One difference exists between these
78
Ti4 Pt3
Ti3 Pt
(a) 1050 ◦C 0h + WQ. 2,000X
Ti4 Pt3
(b) 1050 ◦C 0h + WQ. 5,000X
Figure 4.32: Ti-39Pt sample heat treated at 1050 ◦C followed by an immediate water quenching.
The microstructure reveals that the α-TiPt layer, observed in the periphery of the dendrites in the
as-cast alloy, disappeared and possibly transformed to Ti4 Pt3 .
two experiments, the peak at around 1040 ◦C does not exist in the first heating cycle of the sample
containing 100% α-TiPt in the dendrites, meaning that there is no α-TiPt available to transform
into β-TiPt at the martensitic transformation temperature.
79
0.05
0.03
DSC (mW/mg)
1060 ◦ C
0.2
exo
0.04
1025 ◦ C
DSC (mW/mg)
0.3
exo
670 ◦ C
0.4
heating 1. 40 K/min
heating 1. 20 K/min
cooling 1. 20 K/min
heating 2. 20 K/min
cooling 2. 20 K/min
α-TiPt β-TiPt
0.1
0.0
α-TiPt→Ti4 Pt3
600
700
0.01
0.00
800
900
1000
Temperature (◦ C)
1100
heating 1. 40 K/min
heating 1. 20 K/min
cooling 1. 20 K/min
heating 2. 20 K/min
cooling 2. 20 K/min
−0.01
β-TiPt β-TiPt+Ti4 Pt3
696.8 ◦ C
-12.6 J/g
−0.1
0.02
−0.02
600
1200
(a)
700
800
900
Temperature (◦ C)
1000
1100
(b)
Figure 4.33: DTA results on samples heated following the HT1 treatment. In (a) the initial microstructure contains dendrites of 100% martensite while in (b) the initial microstructure contains
the Ti4 Pt3 +α-TiPt lamellar structure within the dendrites. In both experiments an exothermic peak
is observed around 690 ◦C and is probably associated with the transformation of the α-TiPt phase
to Ti4 Pt3 .
To verify the hypothesis of the transformation of α-TiPt to Ti4 Pt3 at around 690 ◦C, the sample
subjected to the HT1 treatment, containing the lamellar α-TiPt+Ti4 Pt3 structure in the dendrites
(Figure 4.29a), was heat treated at 670 ◦C, 1025 ◦C and 1060 ◦C followed by an immediate water
quench to room temperature after reaching the desired temperature. More Ti4 Pt3 is observed at
670 ◦C (Figure 4.34a) than in the original sample (Figure 4.29a), indicating that must exist a
transformation from α-TiPt to Ti4 Pt3 . This transformation product can be observed in the TEM
image presented in Figure 4.34b where new Ti4 Pt3 (white font) nucleates and grows from the αTiPt.
At 1025 ◦C, almost all the α-TiPt phase had transformed into Ti4 Pt3 (Figure 4.35a). The new
Ti4 Pt3 that nucleates and grows from the martensite possesses many different variants compared to
that which forms from β-TiPt. This can be observed in the TEM image in Figure 4.35b.
80
Ti5 Pt3
α-TiPt+
Ti4 Pt3
Ti4 Pt3
α-TiPt
Ti4 Pt3
Ti3 Pt
(a) 2,000X
(b) 60,000X
Figure 4.34: (a) The microstructure of the sample heat treated at 670 ◦C followed by immediate
quench contains more Ti4 Pt3 within the dendrites. (b) The TEM image shows that new Ti4 Pt3
(white box) is growing on the prior α-TiPt.
Ti5 Pt3
Ti4 Pt3
Ti4 Pt3
Ti3 Pt
(a) 2,000X
(b) 45,000X
Figure 4.35: (a) The microstructure of the sample heat treated at 1025 ◦C followed by immediate
quench reveals that the majority of the α-TiPt transformed into Ti4 Pt3 . (b) This transformation
results in Ti4 Pt3 of different variants within the dendrites.
81
4.4.4
Microstructure of the Ti-44Pt alloys after heat treatment
The Ti-44Pt alloy subjected to the HT1 heat treatment reveals the presence of dendrites of
β-TiPt (which transforms to α-TiPt upon quenching) and Ti5 Pt3 in the interdendritic region. The
absence of Ti3 Pt in the interdendritic regions suggests that oxygen is playing an important role. The
microstructure of the samples subjected to HT2 and HT3 heat treatments, Figures 4.15a and 4.16a
respectively, contains a phase whose contrast was in between Ti5 Pt3 and Ti4 Pt3 which appears as a
medium gray phase in Figure 4.15a. TEM and EDS analyses (Figures 4.15c and 4.15d) revealed
that this phase also corresponds to Ti4 Pt3 , but contains lower platinum (region II) than the Ti4 Pt3
matrix (region I).
In the following, there is an explanation of this phenomenon. According to the DTA scan of
this alloy (Figure 4.1), on cooling at around 1160 ◦C the transformation β-TiPt β-TiPt+Ti4 Pt3
takes place inside the dendrites and the Ti5 Pt3 , existent at 1260 ◦C, remains unchanged to this temperature. An equilibrium is established between Ti5 Pt3 and Ti4 Pt3 (dashed line 1 in Figure 4.36).
At a lower temperature, the Ti5 Pt3 phase begins to dissolve and, therefore, it increases its free energy from 1 to 2 in Figure 4.36. At this lower temperature a new equilibrium between Ti5 Pt3 and
the surrounding Ti4 Pt3 is established, in which Ti4 Pt3 phase with a lower Pt content is transformed
(X1 ).
Ti5 Pt3
G
Ti4 Pt3
2
1
30
36
40
42
X1
44
46
Figure 4.36: Equilibrium condition for the formation of a Ti4 Pt3 of lower Pt content (X1 ) adjacent
to Ti5 Pt3 in the Ti-44Pt alloys subjected to Ht2 and HT3 heat treatments.
82
An important difference between the microstructures observed in the Ti-44Pt alloy subjected
to HT2 and HT3 heat treatments (see Table 4.1 for the nomenclature of the heat treatments) is the
presence of the fine lamellar Ti4 Pt3 +α-TiPt structure in the high Z region of the microstructure
obtained at 800 ◦C (HT3, Figures 4.37a and 4.37b). This transformation seems to be associated
with the additional peaks observed on cooling (between 939 and 1014 ◦C) during the DTA scan on
the Ti-44Pt alloy (Figure 4.1), instead of the formation of oxygen stabilized phases as was proposed
by Biggs et al. [48]. In this case, the β-TiPt phase transforms to the fine lamellar structure on
cooling and, therefore, the martensitic transformation does not take place in this alloy subjected to
HT3. This suggests that a cellular type transformation is taking place in this temperature range,
where β-TiPt decomposes into α-TiPt+Ti4 Pt3 . More investigation is necessary to address this
transformation.
α-TiPt
α-TiPt+
Ti4 Pt3
Ti4 Pt3
Ti4 Pt3
(a) 5,000X
(b) 35,000X
Figure 4.37: Ti-44Pt alloy subjected to HT3 heat treatment. (a-b) A fine lamellar Ti4 Pt3 + α-TiPt
structure transformed in the prior β-TiPt lamellae observed in the microstructure of the Ti-44Pt
alloy subjected to HT2 heat treatment (Figure 4.15a).
83
4.5
Sequence of Transformations and Summary of Observed Phases
The sequence of transformations in the Ti-31Pt alloy is as follows. Above the eutectic tem-
perature, the alloy is comprised of liquid (L) and prior dendrites of Ti3 Pt. At the eutectic temperature, the liquid transforms to eutectic Ti3 Pt+β-TiPt. At the peritectoid temperature (∼1205 ◦C),
Ti3 Pt+β-TiPt transforms to Ti4 Pt3 via the peritectoid transformation. Below the peritectoid, β-TiPt
decomposes into the lamellar Ti4 Pt3 +β-TiPt structure. Finally, the β-TiPt transforms to α-TiPt via
the martensitic transformation. The Ti4 Pt phase is present in the as-cast microstructure because
it is stabilized by C and N (Figure 4.38A). The as-cast microstructure heated to 1260 ◦C (HT1)
is comprised now of Ti3 Pt dendrites and Ti3 Pt+β-TiPt eutectic. The β-TiPt phase transformed to
α-TiPt upon quenching (Figure 4.38B). After the heat treatment HT2, the β-TiPt transforms partitionlessly to Ti4 Pt3 (Figure 4.38C). After the heat treatment HT3, the microstructure obtained
above 1050 ◦C remains unchanged (Figure 4.38C). Ti5 Pt3 is observed after heat treatment because
it is stabilized by oxygen. In the 35Pt alloy, the sequence is very similar with the exception that
during solidification there is no primary phase.
A
as-cast
B
1260 ◦C, 2 h, WQ
C
1050 ◦C, 2 h, WQ
800 ◦C, 2 h, WQ
Ti3 Pt+β-TiPt
Ti3 Pt+β-TiPt
Ti3 Pt+L
Temperature ◦ C
1424 ◦C
Ti3 Pt+β-TiPt
∼1205 ◦C
<1205 ◦C
<1205 ◦C
Ti4 Pt3
Ti4 Pt3 +β-TiPt
1050 ◦C
Ti4 Pt3
∼1000 ◦C
α-TiPt
Final
1260 ◦C
Ti3 Pt+Ti4 Pt3 +
α-TiPt+Ti4 Pt
eutectic transformation
martensitic transformation
α-TiPt
800 ◦C
Ti3 Pt+α-TiPt+Ti5 Pt3
peritectoid transformation
Ti3 Pt+Ti4 Pt3 +Ti5 Pt3
decomposition of β-TiPt
Figure 4.38: Schematic showing the sequence of transformations in the Ti-31Pt alloy in the as-cast
condition (A), after HT1 (B), and after HT2 and HT3 (C). The sequence for the Ti-35Pt is similar
except that in the as-cast condition there is no primary phase.
84
The sequence of transformations in the Ti-39Pt alloy is as follows. Above the eutectic temperature, the alloy is comprised of liquid (L) and prior dendrites of β-TiPt. At the eutectic temperature, the liquid transforms to eutectic Ti3 Pt+β-TiPt. At the peritectoid temperature (∼1205 ◦C),
Ti3 Pt+β-TiPt transforms to Ti4 Pt3 at the peripheries of the dendrites. Below the peritectoid, the
bulk of the β-TiPt dendrites decompose into Ti4 Pt3 in a partitionless manner. Finally, the β-TiPt,
in the eutectic and in the regions near the peripheries of the dendrites, transforms to α-TiPt via the
martensitic transformation (Figure 4.39A). The microstructures after heat treatment are similar to
the ones observed in the Ti-31Pt alloy (Figures 4.39B and Figures 4.39C).
A
as-cast
B
1260 ◦C, 2 h, WQ
C
1050 ◦C, 2 h, WQ
800 ◦C, 2 h, WQ
Ti3 Pt+β-TiPt
Ti3 Pt+β-TiPt
β-TiPt+L
Temperature ◦ C
1424 ◦C
Ti3 Pt+β-TiPt
∼1205 ◦C
<1205 ◦C
Ti4 Pt3
<1205 ◦C
Ti4 Pt3
1050 ◦C
Ti4 Pt3
∼1000 ◦C
α-TiPt
Final
1260 ◦C
Ti3 Pt+Ti4 Pt3 +
α-TiPt
α-TiPt
800 ◦C
Ti3 Pt+α-TiPt+Ti5 Pt3
Ti3 Pt+Ti4 Pt3 +Ti5 Pt3
decomposition of β-TiPt at 1230 ◦C
partitionless tranformation of β-TiPt
condition (A), after HT1 (B), and after HT2 and HT3 (C).
The sequence of transformations for the Ti-44Pt alloy is as follows. Above the eutectic temperature, the alloy is comprised of liquid (L) and dendrites of β-TiPt (larger amount that in the
Ti-39Pt alloy). At the eutectic temperature, the liquid transforms to the eutectic Ti3 Pt+β-TiPt.
Below the peritectoid, the β-TiPt dendrites decompose into Ti4 Pt3 in a partitionless manner at the
interdendritic/β-TiPt interfaces which then grows towards the core of the dendrites. At a lower
temperature, the dendrites decompose into the lamellar Ti4 Pt3 +β-TiPt structure. Finally, the βTiPt, in the lamellar structure and in core of the dendrites, transforms to α-TiPt via the martensitic
85
transformation (Figure 4.40A). Upon heating the as-cast microstructure to 1260 ◦C (HT1), the microstructure is comprised primarily of Ti5 Pt3 in the interdendritic regions and β-TiPt dendrites. After quenching, the dendrites transform to α-TiPt via the martensitic transformation (Figure 4.40B).
After heat treatment HT2, the prior β-TiPt dendrites decompose to the lamellar Ti4 Pt3 +β-TiPt
structure. Upon quenching the β-TiPt transformed to α-TiPt via the martensitic transformation.
The microstructure obtained at 1050 ◦C is furnace cooled to 800 ◦C. Within this temperature
range, the β-TiPt decomposed to Ti4 Pt3 +α-TiPt (Figures 4.39D).
A
as-cast
B
1260 ◦C, 2 h, WQ
C
1050 ◦C, 2 h, WQ
D
800 ◦C, 2 h, WQ
β-TiPt+L
Temperature ◦ C
1424 ◦C
Ti3 Pt+Ti5 Pt3
Ti5 Pt3 +β-TiPt
∼1205 ◦C
1260 ◦C
Ti5 Pt3 +β-TiPt
<1205 ◦C
Ti4 Pt3
Ti4 Pt3 +β-TiPt
<1205 ◦C
1050 ◦C
Ti4 Pt3 +β-TiPt
Ti5 Pt3 +β-TiPt
<1205 ◦C
Ti4 Pt3 +β-TiPt
∼1000 ◦C
α-TiPt
Final
Ti3 Pt+Ti4 Pt3 +
Ti5 Pt3 +α-TiPt
α-TiPt
α-TiPt
800 ◦C
α-TiPt+Ti5 Pt3
α-TiPt+Ti4 Pt3 +Ti5 Pt3
Ti4 Pt3 +α-TiPt
βTiPt+Ti4 Pt3 +Ti5 Pt3
decomposition of β-TiPt at 1230 ◦C
partitionless tranformation of β-TiPt
decomposition of β-TiPt between 1050-800 ◦C
condition (A), after HT1 (B), and after HT2 (C) and HT3 (D).
A summary of the microstructures observed in the as-cast condition and after heat treatment
for Ti-31Pt (Table 4.5), 35Pt (Table 4.5), 39Pt (Tables 4.12, 4.13, 4.14 and 4.15), and 44Pt (Tables 4.16 and 4.17) alloys is presented in the following tables:
86
Table 4.10: Summary of microstructures observed in the Ti-31Pt alloy.
alloy
Condition
Description HT
Microstructure
Ti-31Pt
as-cast
-
Dendrites of primary Ti3 Pt and eutectic in interdendritic regions comprised
of Ti3 Pt, Ti4 Pt3 , α-TiPt and small precipitates of “Ti4 Pt”. The Ti4 Pt3 , αTiPt and “Ti4 Pt” phases were resolved in the TEM (Figures 4.4a and 4.4b).
HT1
1260 ◦C, 2 h WQ
Ti3 Pt dendrites unchanged. Eutectic regions now comprised of Ti3 Pt,
Ti5 Pt3 and β-TiPt (which transforms to α-TiPt upon quenching). See Figure 4.10a.
HT2
1260 ◦C, 2 h, FC1 to 1050 ◦C,
2 h, WQ
Ti5 Pt3 , and Ti4 Pt3 (Figure 4.10b).
HT3
1260 ◦C, 2 h, FC to 1050 ◦C,
2 h, FC to 800 ◦C,2 h, WQ
Ti5 Pt3 , and Ti4 Pt3 (Figure 4.10c).
Ti4 Pt3 +
α-TiPt
Ti3 Pt
BEI as-cast
1
FC=Furnace Cooled
87
α-TiPt
Ti3 Pt
Ti5 Pt3
BEI 1260 ◦C
alloy
Condition
Description HT
Microstructure
Ti-35Pt
as-cast
-
100% eutectic composition. Eutectic now comprised of Ti3 Pt, Ti4 Pt3 and
α-TiPt.
HT1
see Table 4.5
Eutectic now comprised of Ti3 Pt, Ti5 Pt3 and β-TiPt (which transforms to
α-TiPt upon quenching).
HT2
”
Eutectic now comprised of Ti3 Pt, Ti5 Pt3 , and Ti4 Pt3
HT3
”
Eutectic now regions comprised of Ti3 Pt, Ti5 Pt3 , and Ti4 Pt3 .
α-TiPt
Ti3 Pt
Ti5 Pt3
BEI as-cast
88
BEI 1260 ◦C
alloy
Condition
Description HT
Microstructure
Ti-39Pt
as-cast
-
Dendrites of primary β-TiPt that transform partitionlessly to Ti4 Pt3 upon
cooling, a layer of α-TiPt in the peripheries and Ti4 Pt3 peritectoid product observed between the layer of α-TiPt and Ti3 Pt phases. Eutectic in
interdendritic regions comprised of Ti3 Pt, Ti4 Pt3 and α-TiPt.
HT1
see Table 4.5
Prior dendrites of β-TiPt which transforms to α-TiPt upon quenching. Eutectic regions now comprised of Ti3 Pt, Ti5 Pt3 and α-TiPt.
HT2
”
Prior dendrites of β-TiPt transform to Ti4 Pt3 upon cooling from 1260 ◦C to
1050 ◦C. Eutectic regions now comprised of Ti3 Pt, Ti5 Pt3 , and Ti4 Pt3
HT3
”
Prior dendrites that formed Ti4 Pt3 on cooling to 1050 ◦C, remain Ti4 Pt3 .
Prior eutectic regions now comprised of Ti3 Pt, Ti5 Pt3 , and Ti4 Pt3 .
Ti4 Pt3
Ti4 Pt3
α-TiPt
Ti3 Pt
BEI as-cast
89
TEM image as-cast
Table 4.13: Summary of microstructures observed in the Ti-39Pt alloy - Continued.
alloy
Condition
Description HT
Microstructure
α-TiPt+
Ti4 Pt3
Ti3 Pt
α-TiPt
Ti3 Pt
Ti5 Pt3
Ti5 Pt3
BEI 1260 ◦C
BEI 1220 ◦C, WQ
-
1220 ◦C, 1240 ◦C, 1260 ◦C,
WQ
Prior dendrites of β-TiPt decompose to the lamellar Ti4 Pt3 +β-TiPt structure. β-TiPt phase transforms to α-TiPt upon quenching via the martensitic transformation. In addition, Ti4 Pt3 peritectoid product is observed in
the periphery of the dendrites. Eutectic regions now comprised of Ti3 Pt,
Ti5 Pt3 , Ti4 Pt3 , and α-TiPt.
-
1050 ◦C, WQ
Prior dendrites of β-TiPt transform partitionlessly to Ti4 Pt3 upon cooling.
On heating, the layer of α-TiPt observed in the peripheries of the dendrites
disappeared and transformed to Ti4 Pt3 . Eutectic in interdendritic regions
comprised of Ti3 Pt and Ti4 Pt3 .
-
HT1, 670 ◦C, WQ
Prior dendrites of β-TiPt contain the lamellar Ti4 Pt3 +β-TiPt structure
(which transforms to α-TiPt upon quenching). On heating, more Ti4 Pt3
forms on the α-TiPt lamellae (red fonts). Eutectic regions now comprised
of Ti3 Pt, Ti4 Pt3 , Ti5 Pt3 and α-TiPt.
90
alloy
Condition
Description HT
Microstructure
Ti5 Pt3
α-TiPt+
Ti4 Pt3
Ti4 Pt3
α-TiPt
Ti4 Pt3
Ti3 Pt
BEI HT1, 670 ◦C, WQ
-
HT1, 1025 ◦C, WQ
TEM HT1, 670 ◦C, WQ
Prior dendrites of β-TiPt contain the lamellar Ti4 Pt3 +β-TiPt structure
(which transforms to α-TiPt upon quenching). On heating, almost all αTiPt phase transformed to the Ti4 Pt3 phase. The Ti4 Pt3 phase possesses
several different variants (resolved in the TEM). Eutectic regions now
comprised of Ti3 Pt, Ti4 Pt3 , Ti5 Pt3 .
91
alloy
Condition
Description HT
Microstructure
Ti5 Pt3
Ti4 Pt3
Ti4 Pt3
BEI HT1, 1025 ◦C, WQ
HT4
1
1260 ◦C, 2 h, AI2
Ti3 Pt
TEM HT1, 1025 ◦C, WQ
Prior dendrites of β-TiPt decompose to the lamellar Ti4 Pt3 +β-TiPt structure (which transforms to α-TiPt upon cooling). Eutectic regions now
comprised of Ti3 Pt, Ti4 Pt3 , Ti5 Pt3 and α-TiPt.
AI=Air Cooled
92
alloy
Condition
Description HT
Microstructure
Ti-44Pt
as-cast
-
Prior dendrites of β-TiPt followed a sequence of transformations upon
cooling. The peripheries of ∼ 41.5 at.% Pt transform partitionlessly to
Ti4 Pt3 below the peritectoid temperature. Towards the core of the dendrites, the β-TiPt whose Pt content is ∼ 44 at.% Pt decomposed into the
lamellar Ti4 Pt3 +β-TiPt structure. At the martensitic transformation temperature the β-TiPt enriched in Pt (∼ 46 at.% Pt) transforms to α-TiPt via
the martensitic transformation. Divorce eutectic in interdendritic regions
comprised of Ti5 Pt3 and Ti3 Pt phases. Precipitates of Ti3 Pt were observed
in the periphery of the dendrites.
Ti3 Pt
Ti4 Pt3
α-TiPt
Ti5 Pt3
α-TiPt
Ti5 Pt3
BEI as-cast
HT1
see Table 4.5
BEI 1260 ◦C
Prior dendrites of β-TiPt transform to α-TiPt upon quenching. Divorce
eutectic regions now comprised of Ti5 Pt3 phase.
93
alloy
Condition
Description HT
Microstructure
HT2
”
Prior dendrites of β-TiPt transform to the lamellar Ti4 Pt3 (I)+β-TiPt structure upon cooling from 1260 ◦C (coarser morphology). Divorce eutectic
regions now comprised of Ti5 Pt3 and Ti4 Pt3 (II) of a lower Pt content than
Ti4 Pt3 (I). The β-TiPt transforms to α-TiPt upon quenching.
HT3
”
Prior lamellae β-TiPt, contained in the lamellar Ti4 Pt3 (I)+β-TiPt structure, transforms to a fine lamellar Ti4 Pt3 +α-TiPt structure upon cooling
from 1050 ◦C. Divorce eutectic regions comprised of Ti5 Pt3 and Ti4 Pt3 (II)
remains unchanged upon cooling.
Ti4 Pt3 (I)
α-TiPt
Ti5 Pt3
Ti4 Pt3 (II)
BEI 1050 ◦C
94
α-TiPt+
Ti4 Pt3
Ti4 Pt3 (I)
BEI 800 ◦C
CHAPTER 5
CONCLUSIONS AND FUTURE WORK
In this investigation, six Ti-Pt alloys in the composition range 30-50 at.%Pt were studied using
SEM, TEM, and DTA techniques and used to draw the following conclusions.
1. The peritectoid reaction occurs but is limited to the Ti3 Pt/β-TiPt interfaces due to the sluggish nature of the transformation. Therefore, the peritectoid reaction proposed by Biggs and
coworkers [1] was confirmed in this research.
2. A modification to the Ti-Pt phase diagram is proposed taking into consideration the microstructures observed in the as-cast alloys. Specifically, the presence of a layer of α-TiPt
near the periphery of the β-TiPt dendrites, which transform to Ti4 Pt3 upon cooling, in the
Ti-39Pt and Ti-42Pt as-cast alloys (Figures 4.6b and 4.7b) and, an almost constant Pt concentration along the cross-section of the dendrites (Figure 4.18b) suggest that that the β-TiPt
solidus line should be displaced to lower Pt and, also, it should be relatively vertical over
the solidification temperature range of the Ti-39Pt alloy such that little coring accompanies
solidification. Likewise, the β-TiPt solvus must exhibit decreasing Ti solubility (increasing
Pt) until the original solvus line is reached at ∼ 1100 ◦C, which takes into consideration the
composition of the α-TiPt measured at the three different temperatures of the heat treatments
in the Ti-39Pt and Ti-44Pt alloys (Tables 4.8-4.9).
3. The proposed modification of the phase diagram (Figure 4.19) is corroborated by the differences in the volume fraction of dendrites measured experimentally, 29% and the volume
fraction calculated using the Scheil Equation, 50%.
4. The large peak observed around 1230 ◦C in the DTA experiments (Figure 4.1) suggests that
a significant enthalpy is involved in the transformation. The transformation associated with
95
this peak appears to correspond to the decomposition of the β-TiPt phase to the lamellar
Ti4 Pt3 +β-TiPt structure.
5. The Ti5 Pt3 phase was observed in the heat treated alloys and, also, in the Ti-42Pt and Ti44Pt as-cast alloys. It contains a slightly lower platinum content than the Ti4 Pt3 phase.
Considering the information revealed by the CBPs (Figures 3.7a to 3.7c) and the SADPs
(Figures 3.8a to 3.8c) the space group was determined to be P63 /mcm. Based on this space
group, the lattice constants are a ∼ 0.8 nm and c ∼ 0.5 nm.
6. Wavelength Dispersive Spectroscopy (WDS) measurements were carried out on several Ti42Pt heat treated alloys (HT5 and HT6). They reveal that the Ti5 Pt3 phase contains ∼ 35 at.%
Pt and ∼ 5 at.% O, which is considerably larger compared with Ti3 Pt (0.2 at.% O), Ti4 Pt3
(0.7 at.% O) and α-TiPt (0.0 at.% O). This suggests that its formation requires the presence
of oxygen to allow its stabilization in the microstructure. It is likely that the Ti2 Pt(O) phase
observed by Biggs et al. [1] corresponds to the Ti5 Pt3 given the Pt composition measured by
them (33-34 at.% Pt).
7. The Ti4 Pt3 is a highly faulted phase observed in both the hypoeutectic alloys (in the eutectic
regions) and hypereutectic alloys (in both the eutectic and dendrites). WDS measurements
indicated that it has a composition of ∼ 41 at.% Pt. Using a combination of selected area
electron diffraction, convergent beam electron diffraction and careful tilting experiments, it
was determined that it has a complex structure with considerable streaking in the resulting
SADPs (Figure 3.3). The lattice parameters are unknown but must be related to the pseudohexagonal parameters where a ∼ 0.80 nm and c ∼ 2.36 nm.
8. It is speculated that the Ti4 Pt3 phase transforms from a hexagonal phase existent at high
temperatures due to its resemblance to a hexagonal structure and the heavy faulting. To this
point, it is not possible to determine exactly at which temperature the transformation from
this hexagonal phase to the highly faulted phase occurs.
96
9. The Ti4 Pt3 phase can transform from α-TiPt at around 690 ◦C, evolving heat during the transformation. The new Ti4 Pt3 that nucleates and grows from the martensite (α-TiPt) possesses
many different variants compared to that which forms from β-TiPt.
10. A fine lamellar structure of α-TiPt and Ti4 Pt3 was observed in the Ti-44Pt alloy heat treated
at 800 ◦C (HT3). This transformation may be associated with the additional peaks observed
on cooling (between 939 and 1014 ◦C) during the DTA scan of the Ti-44Pt alloy (Figure 4.1).
11. The “Ti4 Pt” phase, observed in the interdendritic regions of the as-cast alloys, is also stabilized by interstitials. WDS measurements revealed that the composition of this phase is
59.6Ti-21.8Pt-15.4N-3.2C. The space group of this phase is Fm3m with a ∼ 0.8 nm.
Considering the findings of this work, there are several observations that need further study.
Up to this point, there is no an explanation for the formation of the fine lamellar α-TiPt+Ti3 Pt
structure in the Ti-44Pt heat treated at 800 ◦C (HT3). There is not conclusive evidence to explain
how, in the Ti-44Pt alloy, it is possible to obtain the α-TiPt without the variants observed in the
alloys heat treated at higher temperatures.
It is possible to refine the calculation of the volume fraction of dendrites to obtain a more
realistic description of the solidification of the alloys instead of using a constant partition coefficient and the Scheil equation. This approach will require the knowledge of the diffusion in the
liquid (or a least an approximate value because incomplete thermodynamic information exists in
the literature for this system). Furthermore, the approach will have to consider that the solidus and
liquidus lines are not straight in the temperature range of solidification.
The “Ti4 Pt” phase possesses a cubic structure with the Fm3m space group and a ∼ 0.8 nm.
This phase is also stabilized by interstitials but does not have the same structure as TiC/TiN (Fm3m
with a ∼ 0.4 nm). Therefore, it will be interesting to determine the atom positions and the number
of atoms required to produce the electron diffraction observed in Chapter 3.
97
It was not possible to determine the crystal structure of the Ti4 Pt3 phase. Nevertheless, the
TEM analyses performed using CBED and SADP revealed that it has a complex crystal structure
with electron diffraction evidence pointing in the direction of a triclinic structure whose lattice
parameters should be related to the hexagonal lattice parameters of the Ti5 Pt3 due to their structural
similarities. To determine the crystal structure, good diffraction patterns (x-ray and/or neutron) are
required, i.e., fine (range 25 to 75 μm) and random powders. Reitveld refinement is necessary to
find the best structure that matches, within certain error, the patterns; if more than one structure
match the pattern, ab-initio calculations may help elucidate the crystal structure and atom positions
that minimize the energy of the system.
98
REFERENCES CITED
[1] T. Biggs, L. A. Cornish, M. J. Witcomb, and M. B. Cortie. Revised phase diagram for the
Pt-Ti system from 30 to 60 at .% platinum. Journal of Alloys and Compounds, 375:120–127,
2004.
[2] H. Nishimura and T. Hiramatsu. On the Corrosion Resistance of Titanium Alloys (2nd Report). The Equilibrium Diagram of the Titanium-Platinum System. The Journal of the Japan
Institute of Metals, 21(7):469–473, 1957.
[3] J. L. Murray. The Pt-Ti (Platinum-Titanium) System. Bulletin Of Alloy Phase Diagrams,
3(3):329–335, 1982.
[4] K. Otsuka and X. Ren. Recent developments in the research of shape memory alloys. Intermetallics, 7:511–528, 1999.
[5] K Otsuka and X Ren. Physical metallurgy of Ti – Ni-based shape memory alloys. Progress
in Materials Science, 50:511–678, 2005.
[6] P. K. Kumar and D. C. Lagoudas. Introduction to Shape Memory Alloys, chapter 1, pages
1–50. Springer US, 2008.
[7] R. D. Noebe, T. Biles, and S. A. Padula II. Advanced Structural Materials: Properties, Design
Optimization, and Applications, chapter NiTi-Base High Temperature Shape Memory Alloys:
Properties, Prospects, and Potential Applications, pages 146–181. CRC, first edition, 2006.
[8] H. C. Donkersloot and J. H. N. Van Vucht. Martensitic Transformations in Gold-Titanium,
Palladium-Titanium and Platinum-Titanium Alloys Near The Equiatomic Composition. Journal of the Less-Common Metals, 20:83–91, 1970.
[9] Grant A. Hudish. Characterization of high temperature shape memory alloys in the Ni-Ti-Pt
ternary system. Master’s thesis, Colorado School of Mines, 2009.
[10] G. R. Purdy. Peritectoid Decomposition, pages 6792–6794. Elsevier, Oxford, 2001.
[11] G. V. Kurjumov and L. G. Khandros. First reports of the thermoelastic behaviour of the
martensitic phase of Au-Cd alloys. Doklady Akademii Nauk SSSR, 66:211–213, 1949.
[12] T. Biggs, M. B. Cortie, M. J. Witcomb, and L. A. Cornish. Platinum Alloys for Shape
Memory Applications. Platinum Metals Review, 47(4):142–156, 2003.
[13] K. Uchino. Shape Memory Materials, chapter Shape Memory Ceramics. Cambridge University Press, 1998.
[14] T. Matsumura, T. Nakamura, M. Tetsuka, K. Takashina, K. Tajima, and Y. Nishi. Shape
Memory Ceramics. In MRS Proceedings, volume 604, pages 161–166, 1999.
99
[15] M. Irie. Shape Memory Materials, chapter Shape Memory Polymer. Cambridge University
Press, 1998.
[16] Andreas Lendlein and Steffen Kelch. Shape-memory polymers. Angewandte Chemie International Edition, 41(12):2034–2057, 2002.
[17] P. Pietrokowski. Novel Ordered Phase, Pt8 Ti. Nature, 206(4981):291–291, 1965.
[18] A. Junod, R. Flukiger, and J. Muller. Supraconductivite et chaleur specifique dans les alliages
A15 a base de titane. Journal of Physics and Chemistry of Solids, 37:27–31, 1976.
[19] M. Li, W. Han, and C. Li. Thermodynamic assessment of the Pt – Ti system. Journal of
Alloys and Compounds, 461:189–194, 2008.
[20] P. J. Meschter and W. L. Worrell. An Investigation of High-Temperature Thermodynamic
Properties in the Pt-Ti System. Metallurgical Transactions A, 7:299–305, 1976.
[21] P. Duwez and C. B. Jordan. The crystal stucture of Ti3 Au and Ti3 Pt. Acta Crystallographica,
5(2):213–214, 1952.
[22] A. E. Dwight, R. A. Conner, Jr, and J. W. Downey. Equiatomic compounds of the transition
and lanthanide elements with Rh, Ir, Ni and Pt. Acta Crystallographica, 18(5):835–839, May
1965.
[23] Reza Abbaschian and Robert E. Reed-Hill. Physical Metallurgy Principles, chapter Deformation Twinning and Martensite Reactions. PWS Publishing Company, Boston, MA, third
edition, 1991.
[24] K. Otsuka and C. M. Wayman. Shape Memory Materials, chapter Introduction. Cambridge
University Press, 1998.
[25] P. A. Stadelmann. JEMS-EMS java version. software, 2004.
[26] Sharvan Kumar, Padam Jain, Seong Woong Kim, Frank Stein, and Martin Palm. An In-Situ
Electron Microscopy Study of Microstructural Evolution in a Co-NbCo2 Binary Alloy. In
Materials Research Society Symposium, volume 1128, 2009.
[27] B. F. Buxton, J. A. Eades, J. W. Steeds, and G. M. Rackham. The symmetry of electron diffraction zone axis patterns. Philosophical Transactions of the Royal Society A,
281(1301):171–194, 1976.
[28] T. K. Biswas and K. Schubert. Einige neue Phasen vom Mn5 Si3 -Typ. Zeitschrift Fur Metallkunde, 58:558–559, 1967.
[29] K. Cenzual and E. Parthé. Zr5 Ir3 with a deformation superstructure of the Mn5 Si3 structure.
Acta Crystallographica, C42:1101–1105, 1986.
[30] Pierre Villars and L. D. Calvert. Pearson’s Handbook of Crystallographic Data for Intermetallic Phases. Volumes 1; 2; 3. THREE VOLUME SET. Asm Intl, 1985.
100
[31] A. Raman and H. Ghassem. Characteristics of the Mn5 Si3 , CaZn5 and CeNi3 Type Phases.
Journal of the Less-Common Metals, 30:185–197, 1973.
[32] A. Iandelli and A. Palenzona. The Ytterbium-Platinum System. Journal of the Less Common
Metals, 43(1-2):205–209, November 1975.
[33] A. Iandelli and A. Palenzona. On the crystal structures of Yb5Ir3. Journal of Less-Common
Metals, 83:L1–L5, 1982.
[34] Joël Le Roy, Jean-Michel Moreau, and Dominique Paccard. Structures of the Rare-EarthPlatinum Compunds R7 Pt3 , R2 Pt, R5 Pt3 and RPt. Acta Crystallographica, B34:9–13, 1978.
[35] J. Le Roy, J M Moreau, and D Paccard. R5 T3 Compounds (R=Rare Earth T=Rh, Ir) with an
Mn5 Si3 -Type Structure. Structure, 86:63–67, 1982.
[36] H. Schachner, E. Cerwenka, and H. Nowotny. Neue Silizide vom M5 Si3 -Type mit D88 Struktur. Monatshefte für Chemie/Chemical Monthly, 85:245–254, 1954.
[37] H. Nowotny, B. Lux, and H. Kudielka. Das Verhalten metallreicher, hochschmelzender Silizide gegenüber Bor, Kohlenstoff, Stickstoff und Sauerstoff. Monatshefte für
Chemie/Chemical Monthly, 87:447–470, 1956.
[38] E. Parthé and J. T. Norton. Crystal structures of Zr5 Ge3 , Ta5 Ge3 and Cr5 Ge3 . Acta Crystallographica, 11(1):14–17, Jan 1958.
[39] W. Jeitschko, H. Nowotny, and F. Benesovsky. Kohlenstoffhaltige ternäre Verbindungen (V-Ge-C, Nb-Ga-C, Ta-Ga-C, Ta-Ge-C, Cr-Ga-C und Cr-Ge-C). Monatshefte für
Chemie/Chemical Monthly, 94:844–850, 1963.
[40] W. Jeitschko, H. Nowotny, and F. Benesovsky. Verbindungen vom Typ T5 M3 X. Monatshefte
für Chemie/Chemical Monthly, 95:1242–1246, 1964.
¨
[41] W. Rieger, H. Nowotny, and F. Benesovsky. Untersuchungen in Systemen: Ubergangsmetall
(T)-Bor-Aluminium. Monatshefte für Chemie/Chemical Monthly, 95:1417–1423, 1964.
[42] K. J. Leonard, R. Tewari, A. Arya, J. C. Mishurda, G. K. Dey, and V. K. Vasudevan. Decomposition of the βo phase to the P63 /mcm, hP18 structure in Nb-(24-36)Ti-40Al alloys. Acta
Materialia, 57(15):4440–4453, 2009.
[43] J. D. Corbett, E. Garcia, A. M. Guloy, W. Hurng, Y. Kwon, and E. A. Leon-Escamilla.
Widespread Interstitial Chemistry of Mn5 Si3 -Type and Related Phases. Hidden Impurities
and Opportunities. Chemistry of Materials, 10:2824–2836, 1998.
[44] C. G. Wilson, D. K. Thomas, and F. J. Spooner. The crystal structure of Zr4 Al3 . Acta
Crystallographica, 13:56–57, 1960.
[45] H. Okamoto. Desk Edition: Phase Diagram for Binary Alloys. ASM International, 2000.
[46] A Palenzona and A Iandelli. The crystal structure and lattice constants of Re3 Pd4 , Y3 Pd4 and
Th3 Pd4 compounds. Journal of the Less-Common Metals, 34(1):121–124, January 1974.
101
[47] K. Cenzual, L. M. Gelato, M. Penzo, and E. Parthé. Overlooked trigonal symmetry in structures reported with monoclinic centered bravais lattices; trigonal description of Li8 Pb3 , PtTe,
Pt3 Te4 , Pt2 Te3 , LiFe6 Ge4 , LiFe6 Ge5 , CaGa6 Te10 and LA3.266 Mn1.1 S6 . Zeitschrift füur Kristallographie, 193(3-4):217–242, 1990.
[48] T. Biggs, M. B. Cortie, M. J. Witcomb, and L. A. Cornish. Martensitic Transformations, Microstructure, and Mechanical Workability of TiPt. Metallurgical and Materials Transactions
A, 32A(August):1881–1886, 2001.
[49] W.S. Rasband. ImageJ. U. S. National Institutes of Health, Bethesda, MA, USA,
http://imagej.nih.gov/ij/. 1997-2011.
[50] M.D. Abramoff, P.J. Magalhaes, and S.J. Ram. Image Processing with ImageJ. Biophotonics
International, 11(7):36–42, 2004.
102
APPENDIX A
XRD AND NEUTRON DIFFRACTION RESULTS
From Chapter 4, it is clear that the Ti4 Pt3 phase observed first by Biggs and coworkers [1] is a
stable phase that can transform via partitionless transformation from β-TiPt during solidification,
or via the transformation β-TiPt α-TiPt+Ti4 Pt3 during heat treatment at around 1230 ◦C. TEM
analyses revealed that it is a highly faulted phase that has some structural similarities to the hexagonal, similar a parameter and similar [0001] diffraction pattern, but the c parameter is very different
with diffraction evidence that suggests a triclinic structure. To this point, there is no clarity as to
the type of crystal structure and atom positions in this phase. Therefore, an effort was made to
determine the structural details using x-ray and neutron diffraction experiments were carried out.
Only preliminary results are presented here.
A-1
Methodology
Neutron and x-ray diffraction experiments were performed in the Spectrometer for Materi-
als Research at Temperature and Stress (SMARTS) at the Los Alamos Neutron Scattering Center
(LANSCE) and in a Phillips (PANalytical) X’Pert PRO x-ray diffractometer at the Colorado School
of Mines, respectively. A as-cast Ti-39Pt powder sample was scanned in the SMARTS spectrometer at four different temperatures: room temperature (RT), 800 ◦C, 1100 ◦C and 1300 ◦C. A
˚
white beam was used for this neutron experiments. The powder was sampled in the range 0.9-4 A.
Unfortunately, the powder was not sufficiently random and information of scattering from low 2θ
angles was not obtained. Likewise, a sample of as-cast Ti-42Pt powder was scanned in the x-ray
diffractometer. In this case, the powder was prepared using an alumina mortar and pestle and then
sieved to a particle size range 53-63 μm in an effort to obtain a random powder suitable for x-ray
˚ and the starting 2θ angle was 7.2◦ .
experiments. The beam used was CuKα (λ = 1.54056 A)
103
The initial approach is to calculate the theoretical intensities for each phase that could be
present in the powders, namely Ti3 Pt, Ti5 Pt3 , α-TiPt and β-TiPt (see Appendix B for its calculation). In this manner, the peaks that do not match these phases should correspond to the Ti4 Pt3
phase. The intensity of diffraction peaks can be calculated using Equation A-1:
2
I = |F | p
1 + cos2 2θ
sin2 θ cos θ
e−2M
(A-1)
where I is the relative integrated intensity, F is the structure factor, p is the multiplicity factor,
1 + cos2 2θ
which depends on the crystal system ans the specific diffraction planes,
is the Lorentz
sin2 θ cos θ
polarization factor, θ is the Bragg angle and e−2M is the temperature factor which take into consideration the thermal vibration of the atoms about their mean positions. In turn, the structure factor
for x-rays can be calculated using Equation A-2:
Fhkl =
N
fn exp [2πi(hun + kvn + lwn )]
(A-2)
n=1
where fn is the atomic scattering factor, N the total number of atoms in the basis, (h, k, l) the
Miller indices of the hkl reflection, and (u, v, w) the position of the atoms in the unit cell. On the
other hand, the structure factor for neutrons is given by:
Fhkl =
N
b exp [2πi(hun + kvn + lwn )]
(A-3)
n=1
where b is the scattering length on the order of 10−12 cm. The scattering lengths for titanium,
platinum and niobium are −3.37 × 10−13 cm, 9.60 × 10−13 cm and 7.05 × 10−13 cm, respectively. To compare directly the x-ray atomic scattering factor f with the neutron scattering length
b, Equation A-4 is used:
f o = re f
(A-4)
where fo is the atomic scattering factor for comparison, and re is the classical radius of the electron
equal to 2.818×10−13 cm. Some differences to take into consideration between neutron diffraction
104
with x-ray diffraction:
i) The scattering intensity is much greater for x-rays than for neutrons because the source is
much more intense.
ii) Neutrons interact with the nucleus of the atoms and such interaction is of short range (∼ 10−5
˚ which is shorter than interatomic distances (∼ 1 A).
˚ Therefore, the scattering length b
A)
is essentially independent of the scattering vector (∼ sinθ/λ). On the other hand, x-rays
interact with the electron cloud surrounding the nucleus, whose extension is comparable with
the interatomic distances and, therefore, the atomic scattering factor f decreases with the
scattering vector.
iii) Neutron scattering lengths vary irregularly with atomic number Z due to the neutrons/nuclei
interaction, and there is not a systematic variation with Z. Therefore, neutron diffraction can
be use to distinguish elements that are close to one another in the periodic table, where x-rays
and electron diffraction fall short.
iv) Neutrons scatter light elements more effectively than x-rays; therefore, it is easier to detect
light atoms (H, C, N and O) in the presence of heavy atoms.
v) Some scattering lengths are negative (e.g. Ti) which enhances the difference between elements. This can provide an effective diagnostic tool when compare with x-ray data.
A-2
Results
The Ti-39Pt powder sample scanned in the SMARTS spectrometer at LANSCE was encap-
sulated in a Niobium can and positioned horizontally inside a furnace capable of reaching temperatures up to 1500 ◦C. Likewise, a as-cast Ti-42Pt powder sample was scanned in a Phillips
(PANalytical) X’Pert PRO x-ray diffractometer at the Colorado School of Mines. Figure A-1a
shows a partial neutron pattern for both heating and cooling (RT, 800 ◦C, 1100 ◦C and 1300 ◦C),
˚ Should be kept in mind that the powder
corresponding to a range of d-spacings, namely 2.5-4.4 A.
105
for neutron diffraction was not totally random and; therefore, it may not show all the orientations.
However, a number of phases were observed as follows:
˚ and 3.5 A
˚ are consistent with the (200) and (110) reflections from the
i) The peaks around 2.5 A
Ti3 Pt phase.
˚ that is present from RT to 1100 ◦C and then disappears at 1300 ◦C.
ii) There is a peak at ∼ 3.44 A
It is not observed on cooling. This peak is observed in the XRD pattern (Figure A-1b) and
therefore, it may corresponds to the Ti4 Pt3 phase.
˚ seem to coincide with the (010) reflection of the αTiPt phase.
iii) The peaks ∼ 2.8 A
˚ observed at 1100 ◦C(both on heating and cooling) and 1300 ◦C is coniv) The peak at ∼ 3.2 A
sistent with the (100) reflection from the β-TiPt phase.
˚ 2.9 A
˚ and 3.35 A
˚ are not consistent with any of the known phases
v) The peaks at ∼ 2.7 A,
in the alloys studied (dashed boxes in Figure A-1a) and, do not have a counterpart in the
XRD pattern. This suggests that some contamination or reaction with Nb occurred during the
experiment.
Further studies are required to determine the structure of the Ti4 Pt3 phase. Any preparation of
powders for diffraction must consider the effect of not having a random sample and, any possible
interaction with the container used in these experiments.
106
0.4
Ti3 Pt (110)
β-TiPt (100)
α-TiPt (010)
0.6
Ti3 Pt (200)
Normalized Intensity
0.8
Ti4 Pt3
RT heating
800 ◦ C heating
1100 ◦ C heating
1300 ◦ C heating
RT Cooling
800C Cooling
1100C Cooling
Nb can
β TiPt
Ti3 Pt
αTiPt
Ti5 Pt3
1.0
0.2
0.0
2.6
2.8
3.0
3.2
3.4
3.6
˚
dspacing (A)
(a)
1.0
RT
β TiPt
Ti3 Pt
αTiPt
Ti5 Pt3
Ti4 Pt3
0.6
0.4
Ti4 Pt3
Normalized Intensity
0.8
0.2
0.0
2.6
2.8
3.0
3.2
3.4
3.6
˚
dspacing (A)
(b)
Figure A-1: Neutron diffraction pattern at three different temperatures: RT, 800 ◦C and 1100 ◦C in
˚ XRD pattern at RT of a Ti-42Pt as-cast powder sample.
the range 2.5-3.7 A.
107
APPENDIX B
STRUCTURE FACTORS FOR XRD AND NEUTRON DIFFRACTION
The structure factors for Ti3 Pt, Ti5 Pt3 , α-TiPt and β-TiPt are:
1. Ti3 Pt: It has a primitive cubic crystal structure whose space group is Pm3n. The atom
positions in this phase are:
2a
Pt
000
0.5 0.5 0.5
6c
Ti
0.25 0 0.5
0.5 0.25 0
0.75 0.25 0
0.5 0.75 0
0 0.25 0.25
0 0.5 0.75
Then, the structure factor Fhkl is given by:
Fhkl
πi
πi
(h + 2l) + exp
(3h + 2l)
= fP t 1 + exp (πi(h + k + l)) + fT i exp
2
2
πi
πi
πi
(2h + k) + exp
(2h + 3k) + exp
(2k + l)
+ exp
2
2
2
πi
(2k + 3l) .
(A-1)
+ exp
2
2. α-TiPt: It has a primitive orthorhombic structure whose space group is Pmma. The atoms
positions are:
2e
Pt
0.25 0.5 z
0.75 0.5 z
2f
Ti
0.25 0 z
0.75 0 z
The z coordinate for Ti is 0.18 and for Pt is 0.68, both determined by Dwight et al. [22].
Then, the structure factor is given by:
Fhkl
πi
πi
πi
(h + 4zl) + exp
(3h − 4zl) + fT i exp
(h + 2k + 4zl)
= fP t exp
2
2
2
πi
(3h + 2k − 4zkl) .
(A-2)
+ exp
2
108
3. β-TiPt: It has a primitive cubic structure whose space group is Pm3m. The Ti atoms are
located at (000), while Pt atoms are located at (1/2 1/2 1/2). The structure factor in this case
is:
Fhkl = fT i + fP t exp [πi(h + k + l)].
(A-3)
4. Ti5 Pt3 : It has a primitive hexagonal structure whose space group is P63 /mcm. This phase
has structural similarities to phases that have the Mn5 Si3 prototype structure [28] with atoms
located at:
4d
Ti1
1/3 2/3 0
2/3 1/3 1/2
2/3 1/3 0
1/3 2/3 1/2
6g
Ti2
x 0 1/4
0 x 1/4
x x 1/4
x 0 3/4
0 x 3/4
x x 3/4
6g
Pt
x 0 1/4
0 x 1/4
x x 1/4
x 0 3/4
0 x 3/4
x x 3/4
According to Raman and Ghassem [31] for Ti2 x = 0.2439 and for Pt x = 0.6003. Therefore, The structure factor in this case is:
Fhkl
2πi
πi
2πi
(h + 2k) + exp
(4h + 2k + 3l) + exp
(2h + k)
= fT i exp
3
3
3
πi
πi
πi
(2h + 4k + 3l) + exp
(4hk + l) + exp
(4kx + l)
+ exp
3
2
2
πi
πi
πi
(4x(h + k) + l) + exp
(4hx + 3l) + exp
(4kx + 3l)
+ exp
2
2
2
πi
πi
πi
(4x(h + k) + 3l)
+ fP t exp
(4hk + l) + exp
(4kx + l)
+ exp
2
2
2
πi
πi
πi
(4x(h + k) + l) + exp
(4hx + 3l) + exp
(4kx + 3l)
+ exp
2
2
2
πi
(4x(h + k) + 3l)
.
(A-4)
+ exp
2
109

PHASE TRANSFORMATIONS AND EQUILIBRIA OF

Transcription

Similar documents

Frank Holmes: How to Avoid Vulnerability to Volatility

Admiral Rustler Leaflet

Document 6506112

Document 6514623

Document 6500958

MATERIAL SAFETY DATA SHEET CERROSAFE `.

World leading supplier

Carpenter Gas Turbine Alloys for the Power Generation Industry

ChemLite_PMI Sell_Sheet_US_5001663_A4_5001664_RevB

Motorsport Line Card