U.S. patent application number 11/904304 was filed with the patent office on 2012-07-12 for armour steel.
This patent application is currently assigned to UNIVERSITY OF PRETORIA. Invention is credited to Maweja Kasonde, Waldo Edmund Stumpf.
Application Number | 20120174749 11/904304 |
Document ID | / |
Family ID | 41022669 |
Filed Date | 2012-07-12 |
United States Patent
Application |
20120174749 |
Kind Code |
A1 |
Stumpf; Waldo Edmund ; et
al. |
July 12, 2012 |
Armour steel
Abstract
A low-carbon martensitic armour steel comprises at least Fe, C,
Si and Ni and has a ratio of yield strength to ultimate tensile
strength of less than 0.7. The steel includes retained austenite at
a volume fraction of at least 1%. The low-carbon martensitic armour
steel can be prepared by subjecting a steel which comprises at
least Fe, C, Si and Ni and which has a martensite start temperature
of less than 210.degree. C. to an austenisation heat treatment step
at a temperature of at least 800.degree. C., quenching the steel,
and subjecting the steel to a tempering step at a temperature of
less than 300.degree. C.
Inventors: |
Stumpf; Waldo Edmund;
(Pretoria, ZA) ; Kasonde; Maweja; (Pretoria,
ZA) |
Assignee: |
UNIVERSITY OF PRETORIA
|
Family ID: |
41022669 |
Appl. No.: |
11/904304 |
Filed: |
September 25, 2007 |
Current U.S.
Class: |
89/36.02 ;
148/332; 148/663; 89/903 |
Current CPC
Class: |
C21D 6/002 20130101;
C21D 9/42 20130101; C22C 38/04 20130101; C22C 38/42 20130101; C22C
38/50 20130101; C22C 38/46 20130101; F41H 5/045 20130101; C21D
2211/008 20130101; C22C 38/48 20130101; C22C 38/44 20130101; C22C
38/02 20130101 |
Class at
Publication: |
89/36.02 ;
148/332; 148/663; 89/903 |
International
Class: |
F41H 5/02 20060101
F41H005/02; C21D 6/00 20060101 C21D006/00; C22C 38/44 20060101
C22C038/44; C22C 38/42 20060101 C22C038/42 |
Claims
1. A low-carbon martensitic armour steel comprising at least Fe, C,
Si and Ni which has a ratio of yield strength to ultimate tensile
strength of less than 0.7 and which includes retained austenite at
a volume fraction of at least 1%.
2. The armour steel as claimed in claim 1, which has a ratio of
yield strength to ultimate tensile strength of less than or equal
to 0.65.
3. The armour steel as claimed in claim 2, which has a ratio of
yield strength to ultimate tensile strength of less than or equal
to 0.6.
4. The armour steel as claimed in claim 1, which comprises
0.37-0.43 weight % C, 0.4-1.2 weight % Si, and 1.8-4.0 weight % Ni,
with the balance being mostly Fe.
5. The armour steel as claimed in claim 1, which comprises also Mn,
Cr and Mo.
6. The armour steel as claimed in claim 5, which comprises 0.5-2.0
weight % Mn, 0.8-1.5 weight % Cr, and 0.5-0.6 weight % Mo.
7. The armour steel as claimed in claim 1, which has a martensite
start temperature of less than 210.degree. C.
8. The armour steel as claimed in claim 7, which has a plate
thickness of from 4.5 mm to 8 mm.
9. The armour steel as claimed in claim 8, which has a plate
thickness of from 4.5 mm to 6 mm.
10. The armour steel as claimed in claim 1, which has a volume
fraction of retained austenite of less than 7%.
11. The armour steel as claimed in claim 1, which has a
micro-structure in which the martensite is predominantly present as
twinned plate martensite and not lath martensite.
12. The armour steel as claimed in claim 11, in which the twinned
plate martensite is combined with retained austenite in the same
micro-structure.
13. The armour steel as claimed in claim 1, in which any cementite
is predominantly present as dispersed particles and not as coarse
strings.
14. The armour steel as claimed in claim 1, which is in the form of
a plate with a thickness .delta.mm, and which has a Ballistic
Parameter BP of at least 0.01 where BP=volume fraction of retained
austenite/exp(.delta.).
15. A method of preparing a low-carbon martensitic armour steel,
the method including subjecting a steel which comprises at least
Fe, C, Si and Ni and which has a martensite start temperature of
less than 210.degree. C. to an austenisation heat treatment step at
a temperature of at least 800.degree. C.; quenching the steel; and
subjecting the steel to a tempering step at a temperature of less
than 300.degree. C.
16. The method as claimed in claim 15, in which the austenisation
heat treatment step is at a temperature of between 870.degree. C.
and 950.degree. C.
17. The method as claimed in claim 15, in which the steel is
subjected to the austenisation heat treatment step for a period of
between 20 minutes and 60 minutes.
18. The method as claimed in claim 15, in which the tempering step
is at a temperature of between 150.degree. C. and 250.degree.
C.
19. The method as claimed in claim 15, in which the steel is
subjected to the tempering step for a period of between 20 minutes
and 60 minutes.
20. The method as claimed in claim 15, in which the steel is in the
form of a plate with a thickness of from 4.5 mm to 8 mm.
21. The method as claimed in claim 15, in which the steel comprises
0.37-0.43 weight % C, 0.4-1.2 weight % Si, and 1.8-4.0 weight % Ni,
the balance being mostly Fe.
22. The method as claimed in claim 15, in which the steel comprises
also Mn, Cr and Mo.
23. The method as claimed in claim 22, in which the steel comprises
0.5-2.0 weight % Mn, 0.8-1.5 weight % Cr, and 0.5-0.6 weight %
Mo.
24. A low-carbon martensitic armour steel produced by the method as
claimed in claim 1.
Description
[0001] THIS INVENTION relates to armour steel. In particular, the
invention relates to a low-carbon martensitic armour steel, and to
a method of preparing a low-carbon martensitic armour steel.
[0002] Over the last few decades, military and security
specifications have been developed for armour steel plates based on
their predicted behaviour when impacted by high velocity rounds.
Hitherto the hardness and the tensile and impact properties (in
particular, the yield strength, the ultimate tensile strength, the
elongation at room temperature and the transverse Charpy-V impact
energy at -40.degree. C.) were the main design parameters for most
armour steel plates. Current specifications for military and
security applications recommend a minimum Brinell hardness range of
540 to 600 BHN or 55 to 60 Rockwell C, and a minimum yield and
ultimate tensile strength of 1500 MPa and 1700 MPa respectively,
while a minimum Charpy-V impact energy of 13 Joules at -40.degree.
C. and a minimum elongation of 6% (50 mm gauge length) is required
in South Africa. These requirements have been successfully achieved
over many years by a suitable combination of chemical composition
and heat treatment parameters of armour steels for plates generally
thicker than 8.5 mm up to 30 mm.
[0003] The inventors believe that neither a higher hardness nor
higher mechanical properties are exclusive or even reliable
criteria for predicting the ballistic performance of martensitic
armour steels. Instead, the ratio between the yield and the
ultimate tensile strength, rather than the hardness, is believed by
the inventors to be more important in approximating the true
dynamic fracture and spalling strength of the armour steel.
[0004] According to one aspect of the invention, there is provided
a low-carbon martensitic armour steel comprising at least Fe, C, Si
and Ni which has a ratio of yield strength to ultimate tensile
strength of less than 0.7, and which includes retained austenite at
a volume fraction of at least 1%.
[0005] Preferably, the ratio of yield strength to ultimate tensile
strength is less than or equal to 0.65, more preferably less than
or equal to 0.60.
[0006] The armour steel may comprise 0.37-0.43 weight % C, 0.4-1.2
weight % Si, and 1.8-4.0 weight % Ni, with the balance being mostly
Fe. Preferably, the Si content is higher than 0.8 weight %.
[0007] The armour steel may comprise also Mn, Cr and Mo. Thus, the
armour steel may comprise 0.5-2.0 weight % Mn, 0.8-1.5 weight % Cr,
and 0.5-0.6 weight % Mo. Instead of, or in addition to, Mo, the
armour steel may comprise Co and/or W.
[0008] The armour steel may have a martensite start temperature of
less than 210.degree. C., preferably less than 200.degree. C.
[0009] The armour steel may have a plate thickness of from 4.5 mm
to 8 mm, typically between 4.5 mm and 6 mm.
[0010] Typically, the armour steel has a volume fraction of
retained austenite of less than 7%.
[0011] The armour steel may have a micro-structure in which the
martensite is predominantly present as twinned plate martensite and
not lath martensite. Preferably, the twinned plate martensite is
combined with retained austenite in the same micro-structure.
[0012] The armour steel typically includes cementite. Any cementite
which is present is preferably predominantly present as dispersed
particles and not as coarse strings.
[0013] The armour steel is typically in the form of a plate with a
thickness .delta.mm. The armour steel may have a Ballistic
Parameter BP of at least 0.01 where BP=volume fraction of retained
austenite/exp(.delta.).
[0014] The armour steel may have a Charpy-V impact energy of less
than 13 Joules at -40.degree. C. The armour steel may have a
hardness which is less than 600 BHN or less than 640 VHN. The
armour steel may have a yield strength of less than 1700 MPa, e.g.
less than or equal to 1500 MPa. The armour steel may have an
ultimate tensile strength of less than 2000 MPa, e.g. less than
1700 MPa.
[0015] According to another aspect of the invention, there is
provided a method of preparing a low-carbon martensitic armour
steel, the method including
[0016] subjecting a steel which comprises at least Fe, C, Si and Ni
and which has a martensite start temperature of less than
210.degree. C. to an austenisation heat treatment step at a
temperature of at least 800.degree. C.;
[0017] quenching the steel; and
[0018] subjecting the steel to a tempering step at a temperature of
less than 300.degree. C.
[0019] The austenisation heat treatment step may be at a
temperature of between 870.degree. C. and 950.degree. C. The steel
may be subjected to the austenisation heat treatment step for a
period of between 20 minutes and 60 minutes.
[0020] The tempering step may be at a temperature of between
150.degree. C. and 250.degree. C., e.g. between 150.degree. C. and
200.degree. C.
[0021] The steel may be subjected to the tempering step for a
period of between 20 minutes and 60 minutes, e.g. between 20
minutes and 30 minutes.
[0022] Typically, the steel is in the form of a plate with a
thickness of from 4.5 mm to 8 mm.
[0023] The steel may comprise 0.37-0.43 weight % C, 0.4-1.2 weight
% Si, and 1.8-4.0 weight % Ni, the balance being mostly Fe.
Preferably, the steel has a Si content of higher than 0.8 weight
%.
[0024] The steel may comprise also Mn, Cr and Mo. Thus, the steel
may comprise 0.5-2.0 weight % Mn, 0.8-1.5 weight % Cr, and 0.5-0.6
weight % Mo. Instead of, or in addition to, Mo, the steel may
comprise Co and/or W.
[0025] The steel may have a martensite start temperature of less
than 200.degree. C.
[0026] The invention extends to a low-carbon martensitic armour
steel produced by the method as hereinbefore described.
[0027] The invention will now be described in more detail, by way
of the following experimental studies and the drawings, and with
reference to five experimental steels G1A, G1B, G2A, G2B and G3,
and three conventional steels A66, M38 and RL5.
BRIEF DESCRIPTION OF THE DRAWINGS
[0028] FIG. 1 shows contours of constant YS/UTS ratio for the steel
G1A;
[0029] FIG. 2 shows contours of constant tensile strength (in MPa)
of the steel G1A;
[0030] FIG. 3 shows contours of constant Charpy impact energy (CIE
in Joules) at -40.degree. C. for the steel G1A;
[0031] FIG. 4 shows lines of constant Vickers hardness of the steel
G1A;
[0032] FIG. 5 shows three-D surfaces representing the Charpy impact
energies of the experimental steels G1A, G1B, G2A and G3;
[0033] FIG. 6 shows three-D surfaces representing the ultimate
tensile strengths of the steel G1A that passed ballistic testing,
and the steel G2B that failed the ballistic testing;
[0034] FIG. 7 shows three-D representations of the ratios YS/UTS of
the experimental steels showing lower values for the steels G1A and
G1B that passed the ballistic testing, and higher values for the
steels G2B and G3 that failed the ballistic testing;
[0035] FIG. 8 shows secondary electron scanning microscopy of the
shear lips of steel G1B tempered at 300.degree. C., showing the
cavities and decohesion around MnS particles upon Charpy impact
testing at -40.degree. C.;
[0036] FIG. 9 shows secondary electron SEM of fracture surface
showing an elongated plate-like inclusion of manganese sulphide,
observed after the tensile test of steel G1B tempered at
150.degree. C., with large cavities around the inclusions of the
MnS;
[0037] FIG. 10 shows secondary electron SEM of fractured surface of
untempered steel G2A after ballistic testing showing no decohesion
at the interface between matrix and MnS particle;
[0038] FIG. 11 shows secondary electron SEM of fractured surface of
untempered steel G2A after ballistic testing showing striations
that suggest a propagation of the fracture front under variable
tensile stress upon ballistic impact;
[0039] FIG. 12 shows thin foil transmission electron microscopy of
the steel G1A after tempering at 180.degree. C. and before
ballistic testing, with a scale bar length of 500 nm;
[0040] FIG. 13 shows thin foil transmission electron microscopy of
the steel G2A after tempering at 180.degree. C. and before
ballistic testing, with a scale bar length of 500 nm;
[0041] FIG. 14 shows thin foil transmission electron microscopy of
the steel G2B after tempering at 180.degree. C. and before
ballistic testing, with a scale bar length of 500 nm;
[0042] FIG. 15 shows a bright field thin foil transmission electron
micrograph of the steel G1A after tempering at 300.degree. C.;
[0043] FIG. 16 shows a bright field thin foil transmission electron
micrograph of the steel G1B after tempering at 300.degree. C.;
[0044] FIG. 17 shows a bright field thin foil transmission electron
micrograph of the steel G1A (0.21% Si) tempered at 400.degree. C.
showing large strings of cementite; and
[0045] FIG. 18 shows a bright field thin foil transmission electron
micrograph of the steel G1B (1.06% Si) tempered at 400.degree. C.
showing coarse cementite.
STUDY 1
[0046] In a series of experimental tempered martensitic steel
alloys it was observed that for a given chemical composition, the
heat treatment parameters for advanced ballistic performance are
different from those required for higher mechanical properties,
rendering the often specified relationship between mechanical
properties and ballistic performance questionable. Systematic
analysis of the microstructures and the fracture surfaces of
thirteen laboratory melted tempered martensitic armour plate steels
were carried out to understand the improved ballistic performance
of these steels of which the mechanical properties were actually
lower than currently specified for military and security
applications. It was, furthermore, observed that the detrimental
effect of inclusions on ballistic performance depends on the
tempering temperature and on the strain rate.
[0047] In this study a comparison of the effect of manganese
sulphide inclusions in "slower" strain rate Charpy impact tests at
-40.degree. C. and tensile tests at room temperature, and in
"higher" strain rate ballistic tests on the fracture mode of plates
was made to explain the discrepancies between the performance
predictions based on mechanical properties as in many current
design specifications, and the observed ballistic performance for
plates of tempered martensitic steels with thicknesses less than
8.5 mm.
[0048] Localised thin foil transmission electron microscopy (TEM)
of ballistic impacted regions of these steels suggested the
presence of high temperatures during the impact that induced phase
transformations from an initial twinned martensite to austenite and
back to an untwined martensite.
[0049] Materials and Experiments
[0050] Chemical Composition and Manufacturing
[0051] Initially five experimental armour steels, namely steels G1A
through to G3 were subjected to standard ballistic testing and
their performance compared to those of three currently produced and
used armour steels, here named A66, M38 and RL5. Their chemical
compositions are shown in Table 1.
TABLE-US-00001 TABLE 1 Chemical compositions in wt % of the initial
five experimental martensitic armour steels and three current
commercially produced armour steels Steel C Mn P S Si Cu Ni Cr Mo V
Nb Ti N G1A 0.39 1.22 0.008 0.003 0.21 0.10 2.99 1.49 0.5 0.006
0.002 0.003 0.0049 G2A 0.37 1.15 0.015 0.011 1.06 0.14 3.8 0.52
0.43 0.008 0.008 0.007 0.0036 G2B 0.39 0.65 0.017 0.009 0.8 0.23
2.8 0.22 0.24 0.003 0.006 0.01 0.0051 G2A 0.37 0.40 0.016 0.011
0.43 0.33 2.3 0.24 0.3 0.006 0.006 0.009 G3 0.34 0.39 0.019 0.012
0.40 0.32 2.43 0.27 0.37 0.009 0.009 0.008 A66 0.37 0.68 0.003
0.002 0.24 0.005 1.9 0.48 0.32 0.004 0.001 0.003 0.009 M38 0.38
0.55 0.004 0.002 0.77 0.1 1.79 0.14 0.36 0.001 Nil 0.007 RL5 0.32
0.86 0.008 0.002 0.18 0.26 2.8 0.79 0.45 0.005 0.001 0.003
0.009
[0052] The 5 kg vacuum melted alloys were cast into a 45
mm.times.70 mm.times.230 mm mild steel mould, the ingots were
solution treated for one hour at 1100.degree. C. before hot rolling
in four passes of 20% reduction each down to 6.+-.0.2 mm thickness
at a finishing temperature of between 950 and 900.degree. C. and
then air cooled. The plates were then austenitised at 900.degree.
C. for 20 minutes, water quenched and tempered at 180.degree. C.
for one hour. The microstructures and the phases present before
ballistic testing were analysed by thin foil transmission electron
microscopy and X-ray diffraction. The plate's sizes for ballistic
testing were 200 mm to 250 mm wide and 500 to 550 mm in length.
[0053] Ballistic Testing
[0054] The plates of these steels were ballistically tested
according to the NATO and to the South African specifications. The
plates were clamped on one of the 250 mm edges and passed the
ballistic test if they withstood 5.56 mm rounds fired at a measured
striking velocity higher than 940.+-.10 m/s, from a distance of 30
metres and 0.degree. obliquity angle in a ballistic tunnel. The
parameters and effect of each impact were recorded and in general,
five rounds were fired at each plate. No light should be visible
through any of the impacted regions on a plate for a successful
test.
[0055] Results of the Ballistic Testing
[0056] Ballistic Report
[0057] Table 2 shows the measured mechanical properties of the
steels used in characterising the armour materials according to the
current specifications. In addition, the martensite start
temperatures (M.sub.s), the volume fractions of retained austenite
(% RA) and the yield to tensile strength ratios (YS/UTS) of the
plates, were also considered as design parameters.
TABLE-US-00002 TABLE 2 Properties of the plates before ballistic
testing Impact Retained energy Thickness Hardness YS UTS austenite
Elongation at -40.degree. C. Ms [.degree. C.] [mm] VHN (30 kg)
[MPa] [MPa] YS/UTS [%] [%] [Joules] measured G1A 6.0 .+-. 0.2 578
880 1650 0.53 6 4 10 196 G1B 6.0 .+-. 0.2 565 960 1700 0.56 4 6 13
210 G2A 6.0 .+-. 0.2 610 1500 2200 0.68 0.6 8 14 255 G2B 6.0 .+-.
0.2 520 1500 2000 0.75 0.6 12 17 271 G3 6.0 .+-. 0.2 490 1300 1700
0.76 0.5 14 18 309 A66 6.5 .+-. 0.2 640 1300 1900 0.68 0.6 6 16
(full 243 size) M38 8.5 .+-. 0.2 620 1300 1800 0.72 0.6 7 16 (full
243 size) RL5 12 .+-. 0.2 540 1400 1700 0.82 <0.5 6 18 (full 285
size)
[0058] The Charpy impact energy was measured at -40.degree. C. on
sub-sized specimens of 5.times.10.times.55 mm due to the thickness
limitation of the heat treated plates. The Charpy impact energies
of the armour steels A66, M38 and RL5 are those specified for full
size specimens by the respective manufacturers.
[0059] The results on the ballistic performance of these eight
martensitic steels are presented in Table 3.
TABLE-US-00003 TABLE 3 Results of the ballistic testing Firing
Steel Distance Ballistic designation [m] performance G1A (6.0 mm)
30 Passed G1A (6.0 mm) 10 Passed G1B (6.0 mm) 30 Passed G2A* (6.0
mm) 30 3 out of 5 failed G2B** (6.0 mm) 30 Failed G3** (6.0 mm) 30
Failed A66 (6.5 mm) 30 Passed M38 (8.5 mm) 30 Passed RL5** (12 mm)
30 Failed
[0060] In the rest of this patent specification, the steels will be
designated with (i) no asterisk=passed (ii) one asterisk=partially
failed and (iii) two asterisks=failed the ballistic test.
[0061] Industrial experience shows that plates of steels A66 and
M38 are not reliable for thicknesses smaller than indicated in
Table 3.
[0062] From the ballistic report it appears that thinner
martensitic plates having lower values of the ratio YS/UTS are more
efficient in resisting ballistic perforation. It, however, also
appears that the higher hardness, higher tensile strength and
ultimate tensile strength are not necessarily appropriate in
indicating good ballistic performance for these martensitic
steels.
[0063] Measured Mechanical Properties
[0064] Sample Preparation and Measurement of the Mechanical
Properties
[0065] The tensile and the sub-sized Charpy specimens of the steels
were austenitised for 30 minutes at 800.degree. C., 850.degree. C.,
900.degree. C. and 950.degree. C. respectively in an Argon
atmosphere and were then water quenched to form the martensitic
microstructure. Tempering treatments of the specimens were carried
out at 150, 180, 200, 250, 300, 350 and 400.degree. C. for
different times between 15 and 60 minutes.
[0066] The tensile specimens were wire cut parallel to the rolling
direction and the yield strength, the ultimate tensile strength and
the elongation were determined using an INSTRON 8500 hydraulic
tensile testing machine. The ultimate tensile strength at room
temperature is considered as the first constraint on the ballistic
performance according to some current design procedures.
[0067] As in the case of the tensile properties, current
specifications exist for the Charpy impact energy of armoured plate
steels, a test that is at a relatively "slower" strain rate than
that during ballistic testing. The Charpy impact energy of the
sub-sized specimens was measured at -40.degree. C. to construct the
second constraint on the ballistic performance.
[0068] The dependence of the mechanical properties and ballistic
performance of these armour plate steels on the M.sub.s temperature
was also examined, as measured in a THETA 734 Single Silica Push
Rod LVDT dilatometer. The thin foils for Transmission Electron
Microscopy were prepared from the 3 mm diameter discs
electro-tube-cut from the armour plates in the quenched and in the
quenched and tempered conditions before and after ballistic
impact.
[0069] The fracture surfaces after tensile, Charpy impact testing,
and after ballistic testing were analysed in secondary electron
mode on a JEOL JSM-6300 scanning electron microscope to determine
the mode of cleavage and the possible role of inclusions in the
fracture mechanism.
[0070] Ultimate tensile strength, yield strength to ultimate
tensile strength ratio and Charpy Impact Energy
[0071] The ultimate tensile strength (UTS), the yield strength to
ultimate tensile strength ratio YS/UTS and the Charpy impact energy
at -40.degree. C. (CIE) of the five steels G1A through to G3 were
measured for different austenitisation and tempering temperatures.
The equation describing the variations of the these three
properties with respect to the heat treatment temperatures were
determined for each of the five steels by surface fitting to the
measured values.
[0072] It was observed that third degree polynomials fitted the
results within the experimental ranges of the austenitisation and
the tempering temperatures. This led to the general mathematical
expression of the curve-fitted surfaces using normalising functions
for the two temperatures as defined in equation (1) for the
tempering temperature and in equation (2) for the austenitisation
temperature. Normalised temperatures were used to minimise the
rounding errors in determining the fitting equations.
T tn = ( T t - T tm ) ( T tm - 25 ) ( 1 ) T an = ( T a - T am ) ( T
am - 800 ) ( 2 ) ##EQU00001##
[0073] where T.sub.t is the actual tempering temperature in degrees
Celsius, T.sub.tm is a mean tempering temperature calculated as
T.sub.tm=(25+300)/2=162.5.degree. C. T.sub.a is the actual
austenitisation temperature in degrees Celsius and T.sub.am is the
mean austenitisation temperature calculated as
T.sub.am=(800+950)/2=875.degree. C.
[0074] The correspondence between normalised and actual values of
the austenitisation and tempering temperatures is given in Table
4.
TABLE-US-00004 TABLE 4 Correspondence between normalised and actual
temperatures in degrees Celsius Normalised temperature -1 -0.75
-0.5 -0.25 0 0.25 0.5 0.75 1 Aus- 800 819 837 856 875 894 913 931
950 teniti- sation temper- ature Tem- 25 60 94 128 163 197 231 266
300 pering tem- perature
[0075] The particular mechanical property (MP) is then fitted by
the equation:
MP(T.sub.an,T.sub.tn)=a(T.sub.an).times.T.sub.tn.sup.3+b(T.sub.an).times-
.T.sub.tn.sup.2+c(T.sub.an).times.T.sub.tn+d (3)
[0076] where the fitting parameters a, b, c and d are polynomials
in T.sub.an and are of the general form:
p=A.times.T.sub.an.sup.3+B.times.T.sub.an.sup.2+C.times.T.sub.an+D
(4)
[0077] A, B, C and D are constants.
[0078] Combining equations (3) and (4) gives a sixth order
non-linear equation for the mechanical property in terms of the
normalised temperatures T.sub.an and T.sub.tn.
MP(T.sub.an,T.sub.tn)=(A.sub.1.times.T.sub.an.sup.3+B.sub.1.times.T.sub.-
an.sup.2+C.sub.1.times.T.sub.an+D).times.T.sub.tn.sup.3+(A.sub.2.times.T.s-
ub.an.sup.3+B.sub.2.times.T.sub.an.sup.2+C.sub.2.times.T.sub.an+D.sub.2).t-
imes.T.sub.tn.sup.2+(A.sub.3.times.T.sub.an.sup.3+B.sub.3.times.T.sub.an.s-
up.2+C.sub.3.times.T.sub.an+D).times.T.sub.tn+(A.sub.4.times.T.sub.an.sup.-
3+B.sub.4.times.T.sub.an.sup.2+C.sub.4.times.T.sub.an+D.sub.4)
(5)
[0079] Mechanical properties of Steel G1A (M.sub.s=196.degree.
C.)
[0080] Fitting Function for the Ultimate Tensile Strength UTS, the
YS/UTS and the CIE
[0081] For illustration purpose the fitted equations of UTS, YS/UTS
and (CIE) at -40.degree. C. of steel G1A are represented
mathematically by the functions (6) to (8):
UTS = ( 94.905 T sn 3 - 51.94 T an 2 - 167.71 T an + 137.43 )
.times. T tn 3 + ( - 177.77 T an 3 - 30.043 T an 2 + 333.18 T an -
302.81 ) .times. T tn 2 + ( - 204.53 T an 3 + 226.44 T an 2 +
114.32 T an - 203.98 ) .times. T tn + ( 268.71 T an 3 + 81.169 T an
2 - 472.81 T an + 1875.9 ) ( 6 ) YS TS = ( - 0.0317 T an 3 + 0.0012
T an 2 + 0.0383 T an + 0.0213 ) .times. T tn 2 + ( - 0.0157 T an 3
+ 0.0282 T an 2 + 0.0158 T an + 0.0352 ) .times. T tn + ( 0.0232 T
an 3 + 0.0209 T an 2 - 0.0584 T an + 0.4607 ) ( 7 ) CIE ( - 40
.degree. C . ) = ( 0.2131 T an 3 - 0.6036 T an 2 + 0.5825 T an +
0.0983 ) .times. T tn 3 + ( 0.3463 T an 3 + 0.6073 T an 2 - 1.0607
T an + 0.5168 ) .times. T tn 2 + ( - 0.8878 T an 3 + 0.6606 T an 2
- 0.6379 T an + 1.9633 ) .times. T an + ( - 0.4262 T an 3 - 0.0011
T an 2 - 0.6425 T an + 6.8781 ) ( 8 ) ##EQU00002##
[0082] The objective function may be written in the form:
YS/UTS.ltoreq.r.sub.0 (9)
[0083] where r.sub.0 is 0.68 for the plates of thicknesses smaller
than 6.5 mm, referring to the ballistic report in Table 2.
[0084] The optimum regions for the properties represented by these
equations are found graphically using two-dimensional projections
of contours of equal height (i.e. iso-lines in the normalised
(T.sub.an,T.sub.tn) planes) as illustrated in FIGS. 1 to 3 for the
steel G1A.
[0085] From FIG. 1, it appears that the YS/UTS ratio increases as
the tempering temperature is increased. Furthermore, it is inferred
from FIG. 2 that a tensile strength higher than 1700 MPa is
obtained for steel G1A when the normalised austenitisation
temperature is lower than -0.1 (T=867.degree. C.) and the tempering
temperature lies between the normalised values of -0.5 and 1, (or
actually between 95.degree. C. and 300.degree. C.). If the
tempering temperature were lower than 95.degree. C., the tensile
strength would become difficult to determine because of the brittle
behaviour of this steel in that condition. From FIG. 3, a Charpy
impact energy at -40.degree. C. that is higher than the specified
13 Joules, is obtained for the normalised temperatures between -1
and 0.2 or lower than an actual 890.degree. C. and the normalised
tempering temperature is above 0.9 or actually above 286.degree. C.
The summary of this discussion is presented in Table 5.
TABLE-US-00005 TABLE 5 Heat treatment conditions predicted to be
favourable for the mechanical properties for steel G1A. Favourable
conditions Austenitisation Property temperature Tempering
temperature Low YS/TS 840.degree. C. to 950.degree. C.
<217.degree. C. High UTS <860.degree. C. 95.degree. C. to
300.degree. C. High CIE(-40.degree. C.) <890.degree. C.
>286.degree. C.
[0086] The optimum heat treatment region of steel G1A for ballistic
application could, therefore, be fixed at austenitisation
temperatures higher than 840.degree. C., and lower than 860.degree.
C. for the high strength cases. However, it is more difficult to
find a compromise between the YS/UTS ratio and the Charpy impact
energy at -40.degree. C. in terms of the tempering temperature.
[0087] The hardness of steel G1A varies with both the
austenitisation temperature and the tempering temperature. The
regression analysis and the surface fitting were developed
following the same procedure as adopted earlier. It was noted that
the hardness of steel G1A decreases very fast to values as low as
450 VHN when the tempering temperature was above 200.degree. C.
That would be very low compared to the value of 600 BHN or 640 VHN
specified for military and security applications in some countries,
e.g. South Africa. The variation of the Vickers hardness with the
austenitisation and the tempering temperatures is then written in
terms of the normalised temperatures as follows:
VH=(28.123T.sub.an.sup.3+46.048T.sub.an-26.241T.sub.an+31.589).times.T.s-
ub.tn.sup.3+(-38.351T.sub.an.sup.3-33.588T.sub.an.sup.2+39.306T.sub.an+0.5-
894).times.T.sub.tn.sup.2+(-24.576T.sub.an.sup.2-1.1265T.sub.an-112.31).ti-
mes.T.sub.tn+(32.687T.sub.an.sup.2+26.939T.sub.an+531.32)
[0088] The lines of constant Vickers hardness of the steel G1A are
shown in FIG. 4.
[0089] The hardness of steel G1A in the quenched condition is
relatively constant when the austenitisation temperature is
increased between 800.degree. C. and 900.degree. C. Above this
austenitisation temperature range, for instance at 950.degree. C.,
however, the maximum hardness that was attained increased. This
increase in hardness is mainly due to two factors; firstly the
solid solution hardening of the parent austenite due to the
increased dissolution of some carbides when the austenitisation
temperature was increased and secondly, to a subsequent decrease in
the M.sub.s temperature leading to a harder untempered martensite
with a greater amount of carbon in solution. At 950.degree. C.
grain growth of the austenite can also become significant. The
decrease in the martensite start temperature may, however, lead to
an increase in the volume fraction of retained austenite and
imposes a limit on the increase of the hardness of the steel. The
grain size of steel G1A after austenitisation for 30 minutes, as
determined by the line intercept method using the line scanning
function of the scanning electron microscope, increased from
7.0.+-.0.8 .mu.m when the austenitisation temperature was
850.degree. C., to 10.+-.0.8 .mu.m when the austenitisation
temperature was 950.degree. C.
[0090] The highest tensile strengths are achieved when the
austenitisation temperature is below 867.degree. C. and it dropped
again above this austenitisation temperature. This effect may also
be related to grain growth and the increase in the volume fraction
of the retained austenite. Therefore, it appears that both the
tensile strength and the hardness increase first with an increase
in the austenitisation temperature, but the upper limit in the
tensile strength occurs earlier than for the hardness.
[0091] The rate of decrease in hardness of steel G1A upon
low-temperature tempering, appears to be slower when the
austenitisation temperature is lower within the range from
800.degree. C. to 900.degree. C. This trend indicates that at
higher austenitisation temperatures the amount of carbon dissolved
in the parent austenite is high, which leads to a higher activity
of carbon in the martensite upon tempering. The sudden change of
slope of the hardness curves in FIG. 4, suggests the existence of
two different mechanisms by which the martensite is softened within
the tempering temperature range used here. The first softening
mechanism is active below 150.degree. C. and the second mechanism,
leading to a sharp drop in hardness, becomes active upon tempering
between 200.degree. C. and 250.degree. C. Tempering this armour
steels between 200.degree. C. and 250 leads to the coarsening of
the metastable transition .epsilon.-carbides or .eta.-carbides
previously formed below 150.degree. C. and to their transformation
into cementite.
[0092] Comparison of Mechanical Properties of Experimental Armour
Steels
[0093] The Charpy impact energies at -40.degree. C., the ultimate
tensile strengths and the ratios YS/UTS of the experimental steels
are graphically represented in FIGS. 5 to 7, for direct
comparison.
[0094] Plates of the steels G1A and G1B passed the ballistic
testing despite their lower Charpy-V impact energies at -40.degree.
C., whereas steel G3** failed.
[0095] The higher strength of the steel G2B**, relatively to steel
G1A, throughout the entire range of heat treatment parameters
applied also is not proportional to the ballistic performance.
[0096] However, it appears from FIG. 7 that the steels G1A and G1B
that passed the ballistic testing are characterised by relatively
lower value of the ratios YS/UTS. FIG. 7 also suggests that lower
YS/UTS ratios may be achieved for the steels G2B** and G3** that
failed the ballistic testing by increasing their austenitisation
temperature to 950.degree. C. and tempering at temperatures lower
than 200.degree. C. Such a heat treatment will increase the volume
fraction of the retained austenite in the martensite, hence
reducing the YS/UTS ratio.
[0097] The ratio YS/UTS of the steel G1B (M.sub.s=210.degree. C.)
remains smaller than 0.68 over the entire range of the
austenitisation and tempering temperatures. The Charpy impact
energy is slightly higher than in the case of steel G1A but it is
lower than the specified 13 Joules within a large range of the heat
treatment parameters and becomes higher than 14 Joules for
tempering temperatures higher than 100.degree. C. before it again
decreases for tempering temperatures higher than 300.degree. C. The
values of the UTS higher than the specified 1700 MPa are also found
within the same range of tempering temperatures before the decrease
below 1600 MPa when the steel G1B is tempered at 300.degree. C.
These variations in the Charpy impact energy and UTS are explained
hereinafter and are attributed to the effect of the manganese
sulphide.
[0098] Lower YS/UTS ratios for steel G2A* (M.sub.s=255.degree. C.)
could be achieved at low normalised tempering temperatures between
-1 and 0, or actually lower than 163.degree. C. This limit is lower
than the 200.degree. C. found in the case of steel G1A. Steels G1A
and G2A* have the same carbon content of 0.39% C but have two
different martensite start temperatures of 196.degree. C. and
255.degree. C. respectively, due to the differences in their
manganese and chromium contents. The tensile strength of steel G2A*
is higher than 1700 MPa throughout the entire range of
austenitisation and tempering temperatures used here. This is very
high compared to steel G1A and is also high compared to the limit
specified for the military and security applications. Steel G2A*
also has tensile elongations larger than 11% when tempered at
200.degree. C. This steel has the highest hardness in the
as-quenched condition and also after tempering below 200.degree. C.
The Vickers hardnesses are above 720 VHN or 640 BHN. Finally it
also has a higher YS/UTS ratio than steel G1A.
[0099] The volume fraction of retained austenite in the steel G2B**
(Ms=271.degree. C.) is lower than the detection limit of about 0.5
volume % that the X-ray diffraction technique could detect.
[0100] Some similarities between steel G2B** and steel G2A* may be
noted. The YS/UTS ratios are higher than in the case of steel G1A
and their tensile strengths are also higher than the specified 1700
MPa throughout the entire range of the austenitisation and
tempering temperatures. Their Charpy impact energies at -40.degree.
C. are also higher than the specified 13 Joules throughout the
entire range of the two heat treatment parameters. Some
resemblances are then expected between the microstructures of these
two steels that differ from the steel G1A. Their martensite start
temperatures are both above 250.degree. C. and no retained
austenite (i.e. <0.5%) was detected by X-ray diffraction. The
ultimate tensile strength of steel G2B** is slightly lower than for
steel G2A* but remains higher than 1700 MPa when the normalised
tempering temperature does not exceed the normalised value of 1 or
actually 300.degree. C.
[0101] The martensite start temperature of steel G3**
(M.sub.s=309.degree. C.) is higher than the martensite start
temperatures of the other four steels. It may be observed that
steel G3** represents the highest values of the YS/UTS ratio of all
of the five steels considered up to here. The YS/UTS ratio of steel
G3** in the untempered condition is in the same range than that of
the tempered steels G1A and G1B. The relatively high values of this
ratio for steel G3** may be due to an auto-tempering effect during
the quenching of this steel, in view of its relatively high M.sub.s
temperature. The Charpy impact energy of steel G3** is higher than
14 Joules and it remains high throughout the entire range of heat
treatment parameters. However the ultimate tensile strength drops
to levels lower than 1700 MPa when the tempering temperature is
higher than 200.degree. C., while the ratio YS/UTS is higher than
0.68.
[0102] General Observations on the Mechanical Properties of Steels
G1A Through to G3**.
[0103] The five experimental armour plate steels considered here
may be classified following their martensite start temperatures,
into three groups as shown in Table 6. The first group comprising
steels G1A and G1B have martensite start temperatures lower than
210.degree. C. The second group comprises steels G2A* and G2B** and
have martensite start temperatures near to 250.degree. C. and the
third group comprises steel G3** which has a martensite start
temperature near to 300.degree. C.
[0104] It may be observed, in FIG. 7, that high martensite start
temperatures lead to high values of the YS/UTS ratio in the
quenched as well as in the tempered conditions. The YS/UTS ratio
increases with an increase in the tempering temperature. The high
values of this ratio with high martensite start temperatures, is
probably a consequence of auto-tempering during quenching. However,
this ratio decreases with an increase in the austenitisation
temperature which leads to grain growth and an increase in the
volume fraction of retained austenite because of the lower
martensite start temperature due to a higher dissolution of
carbides. It appears, therefore, that the volume fraction of
retained austenite in these armour steels becomes the main factor
determining the YS/UTS ratio. The YS/UTS ratio is low with a higher
volume fraction of retained austenite and with low tempering
temperatures. At higher tempering temperatures both retained
austenite and martensite decompose with formation of cementite and
ferrite.
TABLE-US-00006 TABLE 6 Groups of armour steels classified according
to the martensite start temperatures Armour Martensite start Group
steel temperature YS/UTS 1 G1A 196.degree. C. <0.65 G1B
210.degree. C. 2 G2A* 255.degree. C. 0.65 to G2B** 271.degree. C.
0.70 3 G3** 309.degree. C. >0.70
[0105] Armour steels L300 (M.sub.s=285.degree. C.), L500
(M.sub.s=253.degree. C.), A66 (M.sub.s=241.degree. C.) and MR300
(M.sub.s=265.degree. C.) are currently being used in military and
security applications within South Africa. The specifications for
these armour steels are stated in terms of the yield strength that
should be higher than 1500 MPa and the tensile strength, which
should be higher than 1700 MPa. These two strength limits will lead
to values of the YS/UTS ratio close to 0.88 and will lead to the
occurrence of localised yielding during impact. Experience within
the industry has shown that steel A66 has better ballistic
performance than the other three for plate thicknesses between 8.5
mm up to 30 mm. According to the proposed categories according to
M.sub.s temperatures, steels L300, L500 and MR300 may be classified
into the second group, whereas the steel A66 belongs to the
transition between the group 1 and the group 2 of armour steels as
previously defined from their martensite start temperatures.
[0106] The armour steels in group 3 (M.sub.s.about.300.degree. C.)
have intermediate tensile strengths between those in the first and
the second groups. The same observation is valid for their
hardnesses. Hence the second group of armour steels is currently
produced for military applications based on the design philosophy
that would link the expected ballistic performance to the hardness
and the tensile strength of these steels, which does not appear to
be an entirely valid design philosophy.
[0107] The inventors believe that the hardness of armour steel is
not an important determinant. Rather the tensile strength is
considered to be important as it compares well to the true fracture
strength during high-velocity impact. In the present study the
YS/UTS ratio is considered in predicting the ballistic performance
of the armour steels. These three modes of predicting the ballistic
performances using (i) the hardness of the plates, (ii) their
ballistic performance index BPI, or (iii) their YS/UTS ratio are
compared hereinafter.
[0108] Prediction of Ballistic Performances Based on Measured
Mechanical Properties
[0109] The Ballistic Performance Index (BPI).
[0110] The Ballistic Performance Index BPI was introduced by
Srivathsa and Ramakrishnan (B. Srivathsa and N. Ramakrishnan,
Ballistic performance maps for thick metallic armour, Journal of
Materials Processing Technology 96 (1999) 81-91 and B. Srivathsa
and N. Ramakrishnan, A ballistic performance index for thick
metallic armour, Computer Simulation Modelling in Engineering, 3
(1998), pp. 33-40).
[0111] The BPIs of the steels G1A, G1B, G2A* and G2A** calculated
according to the above model are shown in Table 7. For this
calculation a muzzle velocity of 940 m/s was considered. The
average Young's modulus of the steels was assumed to be 200 GPa and
the density 7800 kg/m.sup.3. The reductions in area used in the BPI
and measured by tensile testing, were respectively 6%, 11%, 20% and
8% for these four steels.
TABLE-US-00007 TABLE 7 The ballistic Performance Index of the
experimental steels G2A* G1A G1B (Failed G2B** (Passed) (Passed)
3/5) (Failed) BPI 3.7 3.9 4.6 4.5 YS [MPa] 880 1100 1500 1500 UTS
[MPa] 1780 1897 2200 2000 YS/UTS 0.50 0.58 0.68 0.75
[0112] It is inferred from Table 7 that the BPIs of these steels
are relatively close to each other but with a tendency to predict a
higher ballistic performance for those steels with a higher
strength, which contradicts the experimental observation of the
ballistic test. The BPI is calculated using mechanical properties
measured at lower strain rates and at room temperature. It does not
take into account the temperature rise and its dynamic effect upon
ballistic impact on the microstructure through phase
transformations and transitions in martensitic steels, which, in
turn, affect the localised mechanical properties. This can explain
the apparent contradiction between the BPI and the actual ballistic
performance for these martensitic steels.
[0113] The formula for the BPI, however, has the positive value of
taking into account the effect of the ductility of the steel on its
ballistic performance. It includes the tendency for localised
yielding in steels with a high ductility that leads to poor
ballistic performance. The BPI formula also demonstrates the
decrease in ballistic performance when the velocity of the fired
rounds increases. In the case of these steels, the BPI is
multiplied by a factor of 3 to 4 when the velocity of the round is
reduced from 940 to 400 m/s. But the BPI still predicts a higher
ballistic performance for steels that have higher strengths (at
least in the case of martensitic steels), which is not necessarily
reflected in their actual ballistic performance.
[0114] The assessment of performance by ballistic testing remains
indispensable and confirms the current observation that a clear
relationship between the mechanical properties and the ballistic
performance is still lacking. The Ballistic Performance Index
should then possibly be considered as a qualitative indication of
ballistic performance and may be used for the comparative selection
between different armour materials only when their BPIs are
different by more than a given margin or a ratio yet to be
determined.
[0115] The Specifications for South Africa
[0116] The current specifications for military and security
applications of armour steels in South Africa are: [0117] the
hardness is the main factor determining the ballistic performance
and should be higher than 600 BHN, that is equivalent to 640 VHN;
[0118] the transverse Charpy impact energy of the full size
specimen should be higher than 13 Joules at -40.degree. C.; [0119]
the yield strength of the steel should be higher than 1700 MPa;
[0120] the ultimate tensile strength should be higher than 2000
MPa; and [0121] the minimum elongation on a 50 mm gauge length is
fixed at 6%.
[0122] According to this specification the prediction of the
ballistic performance was favourable for the steel G2A* only, a
steel that resisted two fired rounds but the other three perforated
the plate. On the other hand the plates of the steels G1A and G1B
passed the ballistic test despite their lower hardnesses and
tensile properties than specified. The steel G2B** satisfied all
the requirements of this specification except the hardness, and it
failed the ballistic test. One should conclude then that the high
yield strengths, the high tensile strengths, the high elongations
and the high impact energies of the steels G2A* and G2B** did not
play a decisive role in resisting high velocity impacts.
[0123] The Ratio YS/UTS
[0124] The ratios YS/UTS of the experimental tempered martensitic
steels are presented in Tables 2 and 7 from which it is observed
that the armour steel whose ratios of YS/UTS are lower, had good
ballistic performance in the experimental conditions. However, the
higher UTS, higher YS or Charpy impact energy, considered
separately, appear not to be accurate determinants of the ballistic
performance for martensitic steels.
[0125] It may be observed that steel G3** (group 3) has the highest
impact energy throughout the entire range of the austenitisation
and tempering temperatures, whereas steel G1A and steel G1B (group
1) have the lowest impact energy. Steel G2A* (group 2) has a fairly
intermediate level of Charpy impact energy. It also appears that
the Charpy impact energy of the sub-sized specimens of the armour
steels measured at -40.degree. C., increases when the martensite
start temperature of the armour steel is higher, which is not the
case for the ballistic performance.
[0126] Effect of Tempering Temperature and Inclusions on the
Mechanical Properties and the Ballistic Performance.
[0127] The fracture surfaces after the tensile tests at room
temperature and the Charpy impact tests at -40.degree. C. of the
three groups of armour steels, as classified according to their
martensite start temperatures, were compared by SEM. The effect of
silicon, chromium and manganese contents in their resistance to
low-temperature tempering, was also analysed. The shear lips near
the standard notch of the Charpy specimen as well as the area
situated within the fracture surface at a position below the notch
and near to the area of contact with the striking edge of the
pendulum were analysed by SEM.
[0128] It was observed that tempering at 200.degree. C. improves
the toughness of these armour steels and increases their Charpy
impact energy to above 12 Joules. At the same time the effect of
manganese sulphide particles became significant. The fracture face
of these Charpy specimens became ductile with small dimples formed
near the notch as well as near the impact area away from the notch.
The size of the cavities around the manganese sulphide particles
became larger as illustrated in FIGS. 8 and 9. The decrease in the
Charpy impact energy of the specimens upon tempering above
300.degree. C. may then be partially attributed to the detrimental
effect of the decohesion around the manganese sulphide particles in
a relatively soft martensite when the tempering temperature exceeds
200.degree. C.
[0129] Tempering produces carbides and removes the carbon from
solid solution in the martensite and thus lowers the hardness and
increases the toughness. However the detrimental effect of the
manganese sulphide particles plays a role in the fracture mechanism
of these steels in Charpy impact and tensile tests, and imposes a
limit to their increase in impact energy with fracture cavities of
up to 7 .mu.m that are being formed. The shape of the manganese
sulphide particles has a strong effect on the stress concentration
effect during Charpy impact and tensile testing. The shape is
important but of equal importance here is the very low adhesion
between the ferrite matrix and MnS particles. Specifically, upon
tempering above 250.degree. C. the softening of the martensite
promotes decohesion around the elongated MnS particles. Cavities of
diameters larger than 16 .mu.m were formed upon shearing of the
areas around the MnS particles. The Charpy impact energy of these
armour steels becomes lower once again upon tempering at
400.degree. C.
[0130] Both the cementite and the MnS particles are, therefore,
prejudicial to the resistance against "lower" strain rate impact
loading despite the presence of the soft ferrite. Although not
wishing to be bound by theory, this phenomenon can be explained by
the occurrence of localised high stresses around the elongated
inclusions of the MnS particles. These high stresses are favourable
for the nucleation of voids and their coalescence into cavities
that consequently lead to a decrease of the nominal ultimate
tensile strength of the armour steel upon tempering. The other
reason for this decrease of the ultimate strength is the
decomposition of the martensite itself and the formation of coarse
cementite.
[0131] The sub-sized Charpy specimens of steel G2A* (0.009% S,
0.65% Mn) that have martensite start temperatures near to
250.degree. C., show the same brittle behaviour in the untempered
condition as was the case with the steels G1A and G1B but with a
slightly higher impact energy. Besides the mentioned reasons of the
brittle behaviour in the untempered condition, other inclusions
such as the calcium-aluminium compounds inherited from the casting
process also act as stress raisers and, therefore, may have acted
as crack initiators in the hard untempered martensite during the
tensile test. For steel G2A* also the effect of the MnS inclusions
become observable and large cavities are formed around this type of
inclusion that weakened the armour steel when the tensile or the
Charpy specimens are tempered at temperatures above 200.degree.
C.
[0132] The detrimental effect of the MnS particles in the tempered
martensite was not observed to have the same importance with the
high strain rate ballistic testing. The matrix remains coherent to
the inclusions and no cavities or dimples were perceptible around
the interfaces as illustrated in FIGS. 10 and 11. This fact may
explain the high ballistic performance of steels with relatively
lower tensile strengths and Charpy impact energy values. It is
clear that the effect of the inclusions on the fracture mechanism
depends on the strain rate. At higher strain rates the time is too
short for dimples and cavities to initiate and grow with strain
around the inclusions or other discontinuities. Rather, grain
boundary fracture, with some grains being pulled out of the matrix,
was observed to be the failure mode under ballistic impact.
[0133] It is inferred from FIGS. 8 to 11 that the fracture modes at
"lower" strain rate and at "higher" strain rate (ballistic impact)
are different. This may explain the inadequacies between mechanical
properties measured by means of tensile or Charpy impact tests and
the ballistic performance.
[0134] Study 1 Conclusions The brittle fracture of steels G1A and
G1B indicates that they cannot be used in the untempered condition
because of the risk of spalling if impacted by high velocity
projectiles. The tempering treatment at temperatures ranging
between 150.degree. C. and 250.degree. C. improves the ductility of
the armour steels of Group 1 and 2 at room temperature and at
-40.degree. C. SEM also shows that the tempering treatment enhances
the negative effect of the MnS particles at "lower" strain rates as
in Charpy testing and tensile test. The notch in Charpy testing
enhances the brittle behaviour and intergranular fracture of the
untempered armour steels in Group 1. All potential stress raisers
should therefore be avoided in the manufacture of armour
plates.
[0135] Inclusions had pronounced negative effects on both the
strength and the toughness at "lower" strain rates during tensile
and Charpy impact tests than during the high strain rate ballistic
impact. The softening of the martensitic matrix by tempering at
temperatures above 250.degree. C. enhances the decohesion of the
matrix at the inclusion interfaces.
[0136] The effect of the inclusions on the fracture mechanism
depends on the strain rate. Moreover, phase transformations were
observed in the ballistic impacted regions that were not present in
tensile or Charpy specimens.
[0137] Fracture mechanisms of martensitic steels change drastically
between the "lower" strain rate testing and ballistic impact
testing. This may explain the lack of correlation between
mechanical properties and ballistic performance noticed in the
literature. The yield strength, the ultimate tensile strength and
the elongation measured at room temperature using "slower" strain
rates, and the Charpy impact energy measured at -40.degree. C. are,
therefore, not appropriate for the prediction of ballistic
performance.
[0138] Lower values of the YS/UTS ratio that indicate enhanced
resistance to localised yielding, hence higher ballistic
performance, are generally obtained with austenitising and
tempering temperatures that differ from those necessary for the
achievement of high strengths. Strength based design specifications
for ballistic performance, are, therefore, not necessarily
appropriate on their own and microstructural aspects need to be
introduced to improve the prediction of ballistic performance in
tempered martensitic armour plate steels.
STUDY 2
[0139] In this study an alternative design methodology for tempered
martensitic armour steels is proposed which is based on the effect
of retained austenite on the ratio of the yield to ultimate tensile
strength (YS/UTS), the microstructure of the tempered martensite
and its martensite start temperature M.sub.s. This approach was
developed using 6 mm thick armour plates and later was successfully
applied to the design of eight experimental armour steels with
plate thicknesses ranging from 4.5 to 5.2 mm and tested by a
standard 5.56 mm rounds ballistic test.
[0140] The effect of the hardness on the ballistic performance of
armour steels is not uniquely defined. It depends on the strain
rate (function of the striking velocity) and the thickness of the
plate. It appears that in thicker plates and under low strain rates
the ballistic affected region is more localised than in thinner
plates. Therefore, the ability to resist ballistic perforation
depends on the hardness in the first case, whereas the ability to
deform plastically in a large volume around the impact region
becomes the determinant in the case of thinner plate's performance
under high strain rates.
[0141] The inventors previously studied the thermal effect and the
subsequent phase transformations and transitions that occur when
high velocity rounds impact on martensitic steel armour plates and
which are not accounted for in the BPI model mentioned
hereinbefore. Transmission electron microscopy of the ballistic
impact regions in eight armour steels suggested that the thermal
effect that accompanies the impact and the subsequent phase
transformations, might absorb a significant part of the kinetic
energy of fired rounds and consequently improve the resistance to
ballistic perforation. It would then become necessary to account
for these phenomena in the models for an improved prediction of the
ballistic performance, allowing the design of higher performance
martensitic armour steels.
[0142] The design of advanced performance martensitic armour steels
from microstructural considerations has the benefit of overcoming
the lack of correlation between high strength and high ballistic
performance as observed by the inventors.
[0143] Materials and Experiments
[0144] Chemical Composition and Manufacturing
[0145] Five experimental armour steels, namely the steels G1A
through to G3 referred to in Study 1 were subjected to standard
ballistic testing and their performance compared to those of three
currently produced and used armour steels, here named A66, M38 and
RL5. Their chemical compositions, casting details, hot rolling
processes, heat treatments and the specifications of the ballistic
testing applied to these steels are described in Study 1.
[0146] Results of the Ballistic Testing
[0147] The Ballistic Parameter BP
[0148] The results on the ballistic performance of these eight
martensitic steels are presented in Table 1 in Study 1.
[0149] An alternative Ballistic Parameter (BP) is proposed to
account for the microstructural and the plate thickness effects in
predicting the ballistic performance, based on earlier experimental
results obtained by the inventors on a further 13 martensitic
armour steels. The BP is defined as follows:
BP = RA ( % ) exp ( .delta. ) ( 10 ) ##EQU00003##
[0150] where RA is the volume fraction of the retained austenite in
the martensitic microstructure and .delta. is the thickness of the
plate in millimetres. The choice of this expression for the BP
parameter is based on the proportional lowering of the yield to
ultimate tensile strength ratio (YS/UTS) by the presence of
retained austenite in the microstructure and on the increase of the
effective penetrating mass when the thickness of the plate
increases because of the direct transmission of the linear momentum
to the cylinder of material ahead of the fired round within the
plate.
[0151] The Ballistic Parameters of the experimental plates of
interest are compared in Table 8.
TABLE-US-00008 TABLE 8 Ballistic parameter, YS/UTS and ballistic
performance Name of the RA steel [% vol] .delta. BP YS/UTS G1A 6.0
.+-. 0.5 6.0 .+-. 0.2 0.015 0.50 G1B 4.0 .+-. 0.5 6.0 .+-. 0.2
0.010 0.58 G2A* 0.6 .+-. 0.5 6.0 .+-. 0.2 0.0015 0.68 G2B** 0.6
.+-. 0.5 6.0 .+-. 0.2 0.0015 0.75
[0152] The 6 mm thick plate of steel G2A* (Group 2) withstood only
two of the five rounds fired and, therefore, constituted the
demarcating case between those that passed (Group 1) and those that
completely failed the ballistic test (Group 3). The Ballistic
Parameter BP limit value of 0.010, corresponding to the plate of
steel G1B, was then used for subsequent predictions of the
ballistic performance of armour plates with thicknesses smaller
than 6 mm and was successfully applied to eight further
experimental advanced performance steel armour plates with
thicknesses ranging from 4.7 mm to 5.2 mm, whilst keeping the
ballistic test conditions unaltered.
[0153] The BP can be considered as an attempt to find a direct
relationship between the microstructure and the ballistic
performance instead of an indirect relationship via the mechanical
properties. It was observed that plates with BP higher than a value
of 0.010 passed the ballistic test, those with a BP near to 0.006
also passed the ballistic test, but dynamic cracks were found to
propagate while those plates with a BP.about.0.003 failed the
ballistic test. It was also observed that the subsequent
plastically deformed area around the impact regions increased as
the BP increased within the limits of the experiments.
[0154] It may also be observed that steels G1A and G1B had values
of the yield to ultimate tensile strength ratio (YS/UTS) close to
0.6. These lower values of this ratio indicate the steel's ability
to resist localised yielding; in other words, it indicates the
ability of the material to dissipate the absorbed kinetic energy
through a larger plastic strain around the impact area. This
property increases the volume of the material interacting with the
fired round, offering better resistance to perforation. The
elongation during uniaxial tensile testing indicates the tendency
for localised yielding of the steel when impacted. The inventors
believe that the elongation during uniaxial tensile testing should
be kept lower than 7%. This observation seems contrary to the
current specification that recommends an elongation higher than 6%
for martensitic armour steel plates thicker than 12 mm.
[0155] Microstructure
[0156] Thin foil transmission electron micrographs of the armour
steels G1A, G1B, G2A* and G2B** before ballistic testing are
compared in FIGS. 12 to 14 after a prior tempering at 180.degree.
C. for one hour. This tempering temperature was found to be the
most optimum for the above three steels (Groups 1 and 2) for
achieving a high ductility at room temperature. Fine elongated
strings of carbides were found in twinned martensite plates of
steels G1A and G1B that were aligned parallel to the martensite
plate interfaces as illustrated in FIG. 12 of these two steels that
later passed the ballistic test. On the other hand, coarse carbides
had precipitated within the martensite laths and on the lath
interfaces of steels G2A* and G2B** and these gave poor ballistic
performances.
[0157] It is inferred from FIGS. 12 to 14 that a high ballistic
performance requires a microstructure consisting of twinned
martensite with some retained austenite
without coarse carbides. The precipitation of heavy cementite
should be avoided by controlling the chemical composition, i.e.
through the silicon content of the steel and the tempering
temperature.
[0158] Steels G1A and G1B that gave successful ballistic
performances after quenching and tempering at 150.degree. C. to
250.degree. C. for one hour, were also tempered at higher
temperatures of up to 400.degree. C. to find the upper limit of
tempering before the detrimental coarse cementite starts to make
its appearance. After tempering at 300.degree. C. for 1 hour the
TEM thin foil micrographs in FIGS. 15 and 16, show large strings of
coarse cementite that formed along the plate interfaces of the
steel G1A with 0.21 wt % Si, while in steel G1B with 1.06 wt % Si,
noticeably less of these coarse strings of carbides were formed.
Hence it is likely that steel G1B will still have a satisfactory
ballistic performance, even after tempering at a relatively higher
temperature.
[0159] The retardation in the formation of coarse cementite during
tempering of steel G1B can be attributed to its higher content of
1.06% silicon, an element that is well known for its effect on
delaying the formation of cementite from supersaturated metastable
martensite. Tempering at 400.degree. C. is not acceptable even for
armour steels containing more than 1 wt % silicon because of the
significant softening of the material and the formation of coarse
cementite, as shown in FIGS. 17 and 18.
[0160] With tempering at 400.degree. C., strings of cementite in
the steel G1A (with lower % Si) are formed which coarsened along
parallel directions contrary to the dispersed particles of
cementite that precipitated in the steel G1B with its higher %
Si.
[0161] The Martensite Start Temperature of the Steel
[0162] M.sub.s temperatures of the martensitic armour steels were
estimated, with an absolute error of .+-.15.degree. C., using the
regression Formula (II), which was based on 23 measured values
using dilatometry for armour steels within the range of chemical
compositions (in wt %) of interest here.
M.sub.s(.degree. C.)=548-590C-35Mn-18Ni-14Cr-9.5Mo-12Si (11)
[0163] The effect of the austenitisation temperature on the
martensite start temperature was also analysed for these armour
steels in the range between 800 and 950.degree. C., austenitised
for 10 minutes in each case. The martensite start temperatures of
these armour steels decrease slightly, typically by about 6 to
10.degree. C. when the austenitisation temperature is increased
from 800.degree. C. to 950.degree. C. The increase in the
austenitisation temperature has many consequences, i.e. greater
dissolution of carbides, more solid solution hardening of the
parent austenite as well as austenite grain growth and all of these
can modify the martensitic transformation process. For instance,
greater dissolution of the carbides changes the chemical
composition of the matrix and, hence, the chemical driving force
for the transformation of the austenite into martensite. It also
increases the solid solution hardening of the parent austenite,
which affects the movement of the dislocated transformation front
through the harder austenite.
[0164] The inventors believe that an improved design scheme for
high ballistic performance martensitic armour steels should
consider the following with regard to the chemical composition and
the tempering treatment of the martensitic armour steel:
[0165] a. Carbon is the main alloying element that determines the
hardness of the martensite. A hardness higher than 500 VHN may be
obtained when the carbon content of the armour steel is above 0.37
wt % C.
[0166] b. The silicon content of the steel has a strong effect on
the stability of the martensite upon tempering, as it delays the
softening of the martensite during tempering at higher
temperatures. It also appears to increase the resistance to dynamic
coarsening of the cementite upon ballistic impact.
[0167] Softening of these martensitic steels upon tempering gave
the following threshold temperatures (Table 9) where the hardness
started to drop measurably:
TABLE-US-00009 TABLE 9 Silicon content and the temperature of
softening of the martensite in five experimental armour steels.
Tempering temperature of measurable Steel wt % Si M.sub.s softening
G1A 0.21 196.degree. C. 180.degree. C. G3 0.40 309.degree. C.
200.degree. C. G2B 0.43 271.degree. C. 200.degree. C. G2A 0.8
255.degree. C. 250.degree. C. G1B 1.06 212.degree. C. 300.degree.
C.
[0168] Silicon increases the stability of the martensite by
reducing the chemical activity of carbon and, therefore, becomes
effective in delaying the decomposition of the martensite in steels
in the range between 0.5 wt % to 1.0% C. However, from Table 9 it
appears that the stability of the martensite upon tempering does
not correlate uniquely with the M.sub.s of the steel.
[0169] c. The martensite start temperature of these armour steels
may be approximated with acceptable accuracy using the
experimentally determined regression formula (II).
[0170] d. Higher volume fractions of retained austenite in twinned
plate martensite lead to lower values of the YS/UTS and to an
improved ballistic performance.
[0171] A microstructure-based design procedure for advanced
performance tempered low-carbon martensitic armour steels,
particularly of plates with a thickness of 8.5 mm or less (Group 1)
and for high strength components (Group 2) should consist of the
following five steps:
[0172] Step 1: the Chemical Composition
[0173] Select the chemical composition (in wt %) of the steel
within the following range:
[0174] 0.37-0.43% C, 0.5-2.0% Mn, 0.4-1.2% Si, 0.8-1.5% Cr,
0.5-0.6% Mo, 1.8-4.0% Ni with the final chemical composition chosen
such that the martensite start temperature corresponds to the
following classification and probable areas of application:
TABLE-US-00010 TABLE 10 Classification of martensitic armour steels
Thickness of armour M.sub.s plates Application Group1
<210.degree. C. 4.5-6 mm High ballistic performance Group 2
210.degree. C. < 8.5-20 mm High strength and M.sub.s <
260.degree. C. medium ballistic performance Group 3 >260.degree.
C. 30 mm High toughness and poor ballistic performance
[0175] The M.sub.s temperature of the steel may initially be
approximated using formula (11), during the design process but
later a measured value should be obtained by dilatometry
analysis.
[0176] Step 2: the Heat Treatment
[0177] The austenitisation temperature and the tempering
temperature will determine the final properties of the martensitic
steels. The austenitisation conditions have an effect on the
M.sub.s temperature, hence on the thermodynamics and kinetics of
the martensite transformation, and on the microstructure.
TABLE-US-00011 TABLE 11 Optimum heat treatment and microstructures
High performance armour High strength plates components Category
Group 1 Groups 2 Austenitisation 870 to 950.degree. C. for 20 to 60
800 to 860.degree. C. for 20 to minutes 60 minutes Quenching Water
at room temperature Water or oil medium Tempering 150 to
250.degree. C. for 20 250 to 300.degree. C. for 20 to minutes at
least 60 minutes Retained 1-7% volume fraction Not necessary
austenite YS/UTS <0.6 >0.75 Microstructure Twinned plate
martensite Butterfly and lath with nodular retained martensite
without austenite retained austenite Detrimental coarse cementite
Elongated manganese particles sulphide, other inclusions and
cementite when the tempering temperature increases Strain rate Very
high low
[0178] Step 3: Prediction of the Ballistic Performance
[0179] Perform a phase analysis by XRD, determine the volume
fraction of retained austenite and calculate the BP of the plate.
The minimum required BP should be equal to 0.010.
[0180] Step 4: Assessment of the Performance
[0181] Perform a standard ballistic test to confirm the performance
prediction on armour plate having BP values higher than 0.010.
[0182] Step 5: Microstructure Analysis
[0183] Perform TEM and XRD analysis to confirm the dependence of
the ballistic performance on the microstructure and phases
present.
[0184] The proposed design scheme was successfully applied to the
design of a further eight experimental martensitic armour steels
with plate thicknesses ranging between 4.7 and 5.2 mm and BP values
ranging between 0.006 and 0.055.
[0185] Study 2 Conclusions
[0186] It is possible to predict the ballistic performance of
martensitic armour steels by considering the microstructure,
morphology and the phases inherited from the combination of
chemical composition and heat treatment. The Ballistic Parameter
BP, which includes the volume fraction of the retained austenite in
these steels and the M.sub.s temperature, which determines its
morphology, can be considered as criteria for the classification of
martensitic armour steels in three application groups.
[0187] The microstructure-based design scheme has the advantage of
addressing a direct relationship between the microstructure and
ballistic performance instead of an uncertain and disproven
indirect relationship through the mechanical properties.
[0188] For a given chemical composition, it is possible to design
for a high ballistic performance or for high strength depending on
the heat treatment parameters; a lower YS/UTS ratio indicates a
higher resistance to localised yielding upon impact, hence an
improved resistance to ballistic perforation for a given
composition of the martensitic steel.
* * * * *